A.T.Shonhiwa 1 Introduction and motivation Chapter 1. Introduction and motivation 1.1 Background and motivation Hard materials find extensive use as wear parts and cutting tools. Aluminum oxide exhibits a hardness of 16 - 18 GPa. This hardness has been reported to reach 24 GPa when submicron sized alumina is used 1 . A number of patents exist which disclose the use of alumina as a coating on PcBN, either as a PVD-generated thin coating 2,3,4 or as a polycrystalline coating sintered on top of the PcBN layer at high pressures and high temperatures 5 . Two principal factors have limited the widespread use of alumina tools 6 . These are low fracture toughness and low thermal conductivity , which increases the susceptibility of the tool to damage by thermal shock. Fracture toughness of aluminum oxide based ceramics has been increased through introduction of partially stabilized tetragonal zirconium oxide into the aluminum oxide matrix resulting in a sacrifice in hardness. The addition of transition metal carbide (particulary TiC) has resulted in improved thermal conductivity, hardness and fracture toughness compared to monolithic aluminum oxide 6 . Applications of these two aluminum oxide composites in metal cutting range from cast iron machining at higher cutting speeds to steel at moderate speeds. A number of researchers have also explored the possibility of improving the properties of alumina , either by doping 7 or the addition of strengthening phases 8 . In particular, Nihara 9 has shown that addition
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A.T.Shonhiwa 1
Introduction and motivation
Chapter 1.
Introduction and motivation
1.1 Background and motivation
Hard materials find extensive use as wear parts and cutting tools.
Aluminum oxide exhibits a hardness of 16 - 18 GPa. This hardness has
been reported to reach 24 GPa when submicron sized alumina is used
1. A number of patents exist which disclose the use of alumina as a
coating on PcBN, either as a PVD-generated thin coating 2,3,4 or as a
polycrystalline coating sintered on top of the PcBN layer at high
pressures and high temperatures5.
Two principal factors have limited the widespread use of alumina tools
6. These are low fracture toughness and low thermal conductivity ,
which increases the susceptibility of the tool to damage by thermal
shock. Fracture toughness of aluminum oxide based ceramics has
been increased through introduction of partially stabilized tetragonal
zirconium oxide into the aluminum oxide matrix resulting in a sacrifice in
hardness. The addition of transition metal carbide (particulary TiC) has
resulted in improved thermal conductivity, hardness and fracture
toughness compared to monolithic aluminum oxide6. Applications of
these two aluminum oxide composites in metal cutting range from cast
iron machining at higher cutting speeds to steel at moderate speeds.
A number of researchers have also explored the possibility of improving
the properties of alumina , either by doping 7 or the addition of
strengthening phases 8. In particular, Nihara 9 has shown that addition
A.T.Shonhiwa 2
Introduction and motivation
of nanosized silicon carbide in alumina modifies the mode of fracture
from intergranular to transgranular, with a resulting 30% increase in
fracture strength and an increase in slurry erosive wear resistance by a
factor of 3.
Cubic boron nitride is the second hardest material known to man after
diamond. It owes its hardness to the high degree of covalence which is
isoelectronic to that of diamond. The incorporation of cubic boron nitride
as a second phase in an alumina matrix has not yet been reported. It is
envisaged that the incorporation of cubic boron nitride can result in
improved properties of alumina. In particular fracture toughness and
hardness are most likely to improve by incorporation of a harder phase.
From the rule of mixture the predicted hardness of the resulting
composite can be approximated by the equation
cBNOAlComp xHHxH2
)100(
Where Hcomp is the predicted hardness of the composite, HAl2O3 and HcBN
are the hardness values of Al2O3 and cBN respectively and x is the
volume fraction of cBN in the composite. The presence of cBN grains in
an alumina matrix is also expected to deflect cracks resulting in
improved fracture toughness due to internal stresses.
The main problem in sintering alumina with some cubic boron nitride
arises from the fact that
1. Cubic boron nitride riverts to the soft hexagonal allotrope if heat
treated to temperatures around 1400 o C and
A.T.Shonhiwa 3
Introduction and motivation
2. If firing is done in an oxidizing atmosphere, cubic boron nitride
gets oxidized into B2O3 at temperatures around 1000 o C.
Thus the first problem to be addressed in this research project was to
develop a way of co-sintering the composite at temperatures lower than
the hexagonalisation temperature of cubic boron nitride. This was
achieved by using the reaction bonded aluminum oxide technique. This
technique involves attrition milling mixtures of Al and Al2O3 in a ratio of
say, 50:50 by volume followed by firing slowly in an oxidizing
atmosphere so as to oxidize all the aluminum particles into new
alumina. One main advantage of this technique is that the newly formed
Al2O3 bonds with the original Al2O3 to form a coherent alumina monolith
which sinters at temperatures much lower than the conventional
sintering of alumina.
The issue of cubic boron nitride oxidation was addressed by
separating the heat treatment cycle into two regimes. The first heat
treatment cycle meant for oxidizing aluminum into alumina (RBAO) was
done in air at temperatures not exceeding 1000 o C. This was then
followed by sintering in an inert atmosphere (Argon or vacuum) at 1300
o C.
Thus the main aim of this project was to develop a process of co-
sintering cubic boron nitride and alumina to full density without oxidizing
and/or hexagonalising the cubic boron nitride.
A.T.Shonhiwa 4
Introduction and motivation
1.2 Project Overview
This thesis is divided into five chapters as follows.
The first chapter is an introduction on alumina and cubic boron nitride
as ceramic materials their properties and existing products made from
them.
This is followed by chapter 3 which gives details of all chemicals used,
equipment and analytical techniques employed to characterize
materials.
Chapter 4 deals with the experimental results obtained at various
stages from milling of the raw materials through compaction into green
bodies and initial heat treatment up to final sintering. This chapter also
includes an in depth investigation of the oxidation kinetics of aluminum
into alumina in the reaction bonded aluminum oxide process. In
particular the effects of compaction pressure , temperature and
chemical composition and their effects on the oxidation behaviour of
aluminum are discussed. The interaction of Al2O3 and cubic boron
nitride taking into account the relevant phase diagrams and
thermodynamic considerations are also discussed in this chapter.
The final chapter is conclusion and suggested future work.
Literature review
Chapter 2: Literature review
2.1 Introduction
Working definition of ceramic materials includes all inorganic and
non-metallic materials which can be ionic or covalently bonded
and can be crystalline or amorphous and are produced by the
action of heat10 . This definition includes materials not normally
called ceramics, but which have ceramic –type properties,
notably brittleness. Typical examples of advanced ceramics of
practical importance include borides, carbides nitrides, silicides,
carbon and oxides. Most ceramics are characterized by having
high hardness and low fracture toughness. Inorganic
compounds of groups III, IVa, and VIa in the periodic table have
high hardness, low thermal expansion and high thermal
conductivity and excellent oxidation and corrosion resistance,
hence are candidates for speed cutting tool materials. Among
these compounds alumina Al2O3, silicon nitride Si3N4, titanium
carbide TiC, titanium carbonitride Ti(CN), diamond and cubic
boron nitride cBN are actually being used as cutting tools 11 .
The ideal properties which any cutting material should possess in order
to carry out its function are:
a. Hardness. The cutting tool needs to have a hardness value
greater than that of the work piece in order to with stand the
wear action taking place.
b. Toughness. The cutting tool needs to be sufficiently tough so
that it can withstand any interruptions or vibrations occurring
during the machining process.
A.T.Shonhiwa 6
Literature review
c. Hot strength. This is necessary in order to overcome the heat
generated at the cutting tool-work piece interface.
d. Thermal conductivity. This is necessary so that the heat
generated at the cutting tool – work piece interface should be
conducted way.
Hardness and fracture toughness
Hardness is the resistance of a material to indention by another
material and is directly related to the elastic moduli of the material
which in turn is dependent on the nature of the chemical bonding and
crystal structure of the material 12.The rigidity of the crystal lattice and
the inherent strength of the chemical bonds contribute to the hardness
of the material. The hardest known materials diamond and cubic boron
nitride have cubic crystal systems and strong covalent bonds. Typical
hardness values for most commonly used materials are shown in the
table 2.1..
Fracture toughness can be viewed as a measure of the degree of
brittleness of a material13. In general increasing hardness brings with it
a reduction in toughness implying that those materials in the higher
hardness region (Ceramics) are brittle.
A.T.Shonhiwa 7
Literature review
Table 2.1 Typical hardness values of common materials.
Material Hardness (GPa)
Diamond 75
Cubic Boron nitride 45
Boron carbide 30
Silicon Carbide 26
Alumina 21
Tungsten Carbide 19
Zirconium Oxide 15
Hardened steel (65 HRC) 8
Soft Steel (85 HRB) 1.9
Toughening mechanisms in ceramics
Since ceramics are generally brittle approaches for producing strong
ceramics have been directed at enhancing fracture toughness. Much
work has been done in investigating ways of improving fracture
toughness of ceramics. Much improvements in enhancing toughness of
ceramics has been through the control of microstructural
characteristics. Some of the approaches to enhance the toughness of
ceramics include the following.
a. Microcracking.
If microcracks are formed ahead of a propagating crack, they result in
crack branching, which in turn will distribute the strain energy over a
large area resulting in a decrease in stress intensity factor at the
principal crack tip. Crack branching can also lead to enhanced
toughness because the stress required to drive a number of cracks is
more than that for driving one crack14. One good example of microcrack
toughening is Al2O3 toughened with monoclinic ZrO2. Here microcracks
A.T.Shonhiwa 8
Literature review
occur within regions of local residual tension, caused by thermal
expansion mismatch or by transformation.
b. Particle toughening.
Interaction between particles that do not undergo phase transformation
and a crack front can result in toughening due one of the following:
crack bowing between particles, crack deflection at the particles or
crack bridging by ductile particles. Brittle materials containing a second
phase have been found to have higher fracture toughness than those of
homogenous materials and the toughness increases with increase in
volume fraction of dispersed phase and decreases with dispersed
particle size15. The most effective morphology for deflecting crack
propagation has been found to be rod-shaped grains and whiskers16.
Another crack deflection mechanism for toughening ceramics is as
result of the existence of local residual stress in the vicinity of the
dispersed secondary phase. Composites with 25% vol TiC particles in a
matrix of SiC have been found to have a 60% higher fracture
toughness and 40% higher strength than the matrix material alone. The
improved flexural strength and fracture toughness in this system is
thought to result from crack deflection due to residual stress17.
c. Transformation toughening
This involves a phase transformation of second phase particles at the
crack tip with a shear and dilational component, thus reducing the
tensile stress concentration at the crack tip 14. In composites such as
alumina containing partially stabilized zirconia, the volume change
associated with the phase transformation in zirconia (ZrO2 (t) ZrO2(m))
particles is exploited to obtain enhanced toughness. This transformation
A.T.Shonhiwa 9
Literature review
is accompanied with a volume expansion which result in stresses that
tend to closes the crack leading to an increased toughness 18 . Thus
the fracture toughness of Al2O3 based ceramics can considerably be
enhanced by incorporating fine monoclinic ZrO2 particles. A hot pressed
composite containing 15% vol% ZrO2 has a fracture toughness of 10
MPa m ½ 19 19.
d. Whisker or fiber reinforcement
The toughening of ceramics by brittle fibers and/or whiskers occurs
subject to debonding at the interface20. In the absence of debonding,
because the fiber and matrix typically have comparable toughness, the
composite is brittle and satisfies a rule of mixtures (figure 2.1).
Debonding reduces the amplitude of the stress concentration at the
fiber along the matrix crack front and, when sufficiently extensive,
allows the crack to circumvent the fiber, leaving the fiber intact in the
crack wake. The intact fiber inhibits crack opening and allows a
composite toughness exceeding that of either constituent ( figure 2.1).
Figure 2.1 The role of debonding in whisker toughening 20
A.T.Shonhiwa 10
Literature review
Other factors affecting hardness and toughness of ceramic
materials.
The actual hardness of a material depends on several factors, some of
which are grain size, density and purity just to mention a few. Z. Misirli
et al 21 did some work to demonstrate the effect of additives on the
microstructure hardness and fracture toughness of alumina ceramics. In
their work they evaluated hardness and fracture toughness for alumina
samples with various Silica contents. They found out that both hardness
and toughness increase as SiO2 content of alumina decrease. They
attributed this degradation in properties to the increasing amount of
glassy phase at the grain boundaries.
In yet another development A. Krell 22 did some work which showed
that significant increase in hardness can be obtained by reducing the
grain size of sintered alumina down to the submicron range. An
explanation offered for this improvement was the reduction in grain pull
out frequency which is directly related to wear and hardness.
A. Muchtar et al23 have also shown that decreasing grain size can result
in enhancement of fracture toughness of alumina based ceramics. In
this case fracture toughness enhancement was attributed to shift of
fracture mode from trans-granular in coarse grained samples to inter-
granular in submicron grained samples.
A milestone in improving fracture toughness of alumina by decreasing
grain size was reported by R.S. Mishra and A. K. Mukherjee 24 who
proved that the toughness of alumina can be increased to 8MPa m1/2
by using nanosized grains.
A.T.Shonhiwa 11
Literature review
2.2 Alumina
Aluminum oxide (Al2O3) commonly known as alumina is one of the most
widely used, technical ceramic material. Alumina as a raw material
occurs abundantly in nature, most often as impure hydroxides which
are the essential constituents of bauxite ores. Bauxite is an impure
mixture of gibbsite Al(OH)3 = α Al2O3.3H2O, boehmite and diaspore
which are polymorphs of AlO(OH)= Al2O3.H2O respectively 25. The
usefulness of alumina hinges on its properties namely, high melting
temperature, chemical resistance, electrical resistance and hardness 25
.A diverse range of types of alumina exists with a wide range of
properties as shown in table 2.2.. The major markets for alumina-
based ceramics on a weight basis are refractories (50%), abrasives
(20%), whitewares and spark plugs (15%) and engineering ceramics
(10%) 26.
A.T.Shonhiwa 12
Literature review
Table 2.2 Properties of Alumina ceramics 27
Property
Symbol
Units
C610
Mullite
ceramic
50-65%
Al2O3
C620
Mullite
ceramic
65-80%
Al2O3
C780
Aluminum-
oxide
80-86%
Al2O3
C786
Aluminum-
oxide
86-95%
Al2O3
C795
Aluminum-
oxide
95-99%
Al2O3
C799
Aluminum-
oxide
>99%
Al2O3
Density ρ g/cm3 2.6 2.8 3.2 3.4 3.5 3.9-3.98
Strength ζ MPa 120 150 200 250 280 300
Hardness Hv - - - - - 1600-
1700
1800-
2200
Fracture
Toughness
KIC
MPam1/2
- - - - 4 4
Specific
resistivity
@ 20 oC
ρv>20
Ωm
1011 1011 1012 1012 1012 1012
Specific
resistivity
@ 600 oC
ρv>600
Ωm 104 104 105 108 106 108
Thermal
expansion
@30-600
oC
α30-600
10-6K-1
5-7
5-7
6-8
6-8
6-8
7-8
Specific
heat
capacity @
30-600 oC
Cp30-600
JKg-1
K-1
850-
1050
850-
1050
850-
1050
850-
1050
850-
1050
850-
1050
Thermal
conductivity
λ30-100 Wm-1
K-1
2-6 6-15 10-16 14-24 16-28 19-30
A.T.Shonhiwa 13
Literature review
Crystal structure of Alumina and transition aluminas
Aluminum oxide, commonly referred to as alumina possesses strong
ionic inter-atomic bonding giving rise to its desirable material
characteristics. It exists in several crystalline phases which all rivert to
the most stable hexagonal α phase ( corundum) at elevated
temperatures. This is the phase of particular interest for structural/
engineering applications.
Many processes such as the oxidation of Aluminum metal and heating
of gibbsite ores result in the formation of intermediate metastable
alumina phases before the stable α phase is formed25. These
transitional phases are denoted as γ(Gamma), χ(Chi) , η(Eta), ε,
δ(Delta) , θ(Theta)and Ќ(Kappa) and are of importance because of
their use as catalysts or catalyst supports, adsorbents coatings and soft
abrasives. The sequence of transition aluminas that forms is strongly
dependent on the starting material its coarseness and crystallinity,
heating rate, the amount of water vapor in the atmosphere and by
impurities present. The sequences of transition aluminas are given in
figure 2.2.1 28 29 . These sequences are generally accepted, although
there is no clarity on the X-ray identification of some phases and the
existence of others. Transition aluminas have partially disordered
crystal structures all based on a close –packed oxygen sublattice with
varying interstitial aluminum configurations. Transition aluminas can not
be considered true polymorphs of α-alumina. The low temperature ones
in particular may contain some residual OH anions. Moreover, the
sequence of transformation is not reversible, that is, neither α-alumina
nor any of the high temperature aluminas can be converted to one of
the transition aluminas that occur at a lower temperature and therefore
may be classified as thermodynamically unstable. Crystallographic
properties of transition aluminas are given in table 2.3
A.T.Shonhiwa 14
Literature review
Figure 2.2. Phase transformation sequences of aluminum hydroxides 25
Enclosed area indicates range of stability. Open area indicates range of
transition. Path b is favored by moisture, alkalinity and coarse particle
size (100μm): path a by fine crystal size (below 10 μm). As equilibrium
is approached the structures become more ordered forming a
hexagonal oxygen sublattice until stable α- alumina is formed. Unlike
the transition aluminas the crystal structure of α-alumina is well known.
The crystal structure is often described as having O2- anions in an
approximately hexagonal close packed arrangement with Al3+ cations
occupying two-thirds of the octahedral interstices as shown in figure 2.3
crystallographic properties of transitional aluminas are shown in table
2.2.
Bayerite
α
κ
δ θ
η
α θ
Temp o C
a
b
b
η
400 200
γ
θ
1000 600 800
α
1200
Boehmite
Gibbsite χ
a
α θ
Diaspore
α
α
Bayerite
A.T.Shonhiwa 15
Literature review
Figure 2.3 Crystal structure of α-alumina 26
Table 2.2 Crystallographic Parameters of transition aluminas
Phase Crystal
system
Unit Cell Parameters(Angstrons)
a b c
Alpha Cubic 4.98
Chi Cubic 7.95
Eta Cubic 7.90
Gamma Tetragonal 7.95 7.95 7.79
Delta Tetragonal 7.96 7.96 23.47
Iota Orthorhombic 7.73 7.78 2.92
Theta Monoclinic 5.63 2.95 11.68
Kappa Orthorhombic 8.49 12.73 13.39
A.T.Shonhiwa 16
Literature review
Sintering of alumina
Sintering is a process whereby a material , usually in the form of a
powder is subjected to heat treatment resulting in particles bonding
together to form a coherent body with reduced porosity , increased
density and improved hardness, toughness and strength.
From a processing point of view it is important to note that there are
several processing variables that affect densification/ sintering of a
material hence properties of the final product. These include initial
green density of compact, temperature, time, heating rate, particle size
and particle size distribution, purity of starting material and chemical
additives.
Effect of heating rate on the sintering of alumina.
The actual effect of heating rate on densification and grain growth is not
clear, and different research groups have reported conflicting
experimental results. For example Stanciu et al 30 reported that the
final grain size scaled inversely with the heating rate. This is contrary to
Murayama and Shin 31 who reported that the grain growth was
enhanced by faster heating rate. Y. Zhou et al 33 tried to explain this
disparity by using two Al2O3 powders with different particle sizes and
sintering them to different temperatures at different heating rates.
They found out that in general rapid heating rate resulted in reduced
grain growth and the level of reduction depended on the initial powder
size and sintering temperature.
However the effect of heating rate on densification was not monotonic.
In the early stages of sintering, where densification is just starting faster
heating rate resulted in higher densities and at later stages where
densification had proceeded to rather high degrees faster heating rate
led to lower densities.
A.T.Shonhiwa 17
Literature review
Effect of grain size and grain size distribution on the sintering of
alumina
It is well known in classical sintering theory that during sintering
densification and grain growth are two competing processes, and both
of them are driven by the capillary force that is proportional to the
reciprocal of grain size. Thus the smaller the initial powder size, the
larger the densification and grain growth rates during sintering 32. Work
done by Y. Zhou et al 33 has shown that under identical sintering
conditions powders with finer initial particles always attain higher
densities and larger grain growth compared to powder with coarser
particles. In addition powders with finer particles also start to densify at
lower temperatures and densify at greater rates compared to powders
with coarser particles 33 .
The main obstacles in obtaining ceramics with theoretical density have
been attributed to non uniformities in the green bodies and particle
size distribution and degree of agglomeration of the starting powder
are the main origins of the non uniformities 34.
While it is well known that packing of a powder with a bimodal particle
size distribution results in higher green density than a mono sized
powder due to the effective interspace filling between coarse particles
by fine particles 35 enhanced densification in compacts prepared from
powders with bimodal or wide size distribution has not been observed.
On the other hand powders with narrow size distribution have been
reported to result in sintering to high final densities mainly because of
their uniform pore size distribution 36 37 .
In a separate development Tsung-Shou and Michael D. Sacks 38
investigated the effect of grain size distribution on the densification and
sintered microstructure of Al2O3. In their work agglomerate-free
A.T.Shonhiwa 18
Literature review
powders having the same median particle size, but different widths of
distribution , were prepared by sedimentation of high purity commercial
aluminas. Compacts prepared with broad particle size distribution
powder had a higher green density and smaller median pore channel
radius compared to compacts prepared with narrow particle size
distribution powder, indicating that fine particle were efficiently filling the
interstices formed by larger particles. Both narrow and broad size
powders reached final density at the same time/temperature schedule
and had essentially the same average grain size and grain size
distribution. It should be noted however that some experimental
observations 39 suggest that the problem of exaggerated grain growth
may arise if the particle size distribution of the starting powder becomes
too broad.
Effect of impurities on the sintering of alumina.
Research has been done to investigate the influence of minor chemical
constituents on the sintering of alumina40. Previous studies have
already shown that MgO is a beneficial sintering aid while CaO and
SiO2 have deleterious influence on the sintering of
alumina41,42.Previous experimental studies have shown that abnormal
grain growth is strongly related to the presence of impurities, most
notably CaO and SiO2 which form an intergranular liquid phase
(anorthite) which induces grain faceting leading to a more tabular grain
morphology which eventually leads to abnormal grain growth. In their
work S. Bae and S. Baik 40 demonstrated that abnormal grain growth is
not an intrinsic property of commercial alumina but rather is an extrinsic
property controlled by minor constituents that can be present in the
original powder or introduced during powder processing and
A.T.Shonhiwa 19
Literature review
subsequent sintering. In conventional sintering practice the furnace
wall , heating elements are also possible sources of contamination.
Influence of atmosphere on the sintering of alumina
The influence of atmosphere on the sintering behavior of alumina has
been studied extensively. Early studies revealed that the gases in the
sintering atmosphere must be soluble in alumina in order for the near-
theoretical density to be achieved 43 , 44. Insoluble gases generate a
back pressure which opposes the shrinkage of pores and thereby
reduces the driving force for densification. In general solid state
sintering of alumina in reducing atmosphere (hydrogen) results in fully
densified products while products fired in air have residual porosity.
This porosity remains because in the last stages of sintering all of the
pores are isolated within the oxide grains and further shrinkage would
require the pore gas to dissolve in the oxide and diffuse to the external
surface via grain boundaries. Nitrogen is not soluble in alumina at the
sintering temperature and therefore the pores only shrink until the
increased internal gas pressure balances the reduction in surface
energy during the process. Hydrogen on the other hand is soluble and
diffuses rapidly out of the system 26. More recently it has also been
shown 45 that a fully dense hot-pressed alumina will swell if annealed in
an atmosphere containing sufficient quantity of oxygen. In this case the
oxygen reacts with the impurities to produce gases which will then
generate pressures high enough to nucleate voids within the structure.
It has also been demonstrated that the sintering atmosphere affects the
morphological development of the final microstructure. Mocellin et al 46
,47. observed that when alumina was sintered in hydrogen, the pore
A.T.Shonhiwa 20
Literature review
phase was predominantly confined to the grain boundaries , whereas in
nitrogen or oxygen the pores became entrapped within the grains.
Thompson and Harmer 48 investigated the effect of atmosphere on the
final stage sintering kinetics of ultra –pure alumina. In particular they
investigated the effect of oxygen partial pressure on densification rate
and grain growth rate. They concluded that sintering in low oxygen
partial pressure enhances densification rate and increases grain growth
rate. Additionally it was also observed that sintering in low oxygen
partial pressure enhances relative pore mobility and reduces the
susceptibility to pore/ boundary separation.
Pressure assisted sintering.
Sintering with the aid of mechanical pressure is called hot pressing. The
sample is heated to high temperatures and mechanical pressure is
applied to increase the driving force for densification by acting against
the internal pore pressure without increasing the driving force for grain
growth. Practical advantage of hot pressing is that dense samples with
minimal grain growth can be obtained at much lower temperatures49 .
The material to be hot pressed is first precompacted before being
placed in a hot press die to get a reasonable green density. One major
disadvantage of this technique is that sample shapes are limited.
Another pressure assisted sintering technique was described by
Hardtl50 . This technique is called hot isostatic pressing. In this
technique no die is used and an inert gas is employed as an isostatic
pressure medium. Because the working fluid is a gas it is necessary
that the ceramic to be hot pressed be sintered first. This first sintering
must yield a material having no open or interconnected porosity
otherwise no force will be transmitted to the component. One
advantages of this technique over hot pressing are that the sample
A.T.Shonhiwa 21
Literature review
shape is not critical since pressure is applied isostatically. Another
additional benefit is the elimination of unwanted reactions between
sample and die walls which can be a problem in uniaxial hot pressing.
Alumina based composites
Although alumina is one of the widely used technical ceramics because
of its low density, high strength, high hardness and high temperature
capability it however has some draw backs. For an example its
fracture toughness makes it difficult to withstand severe conditions
applied for example, in the field of high-speed cutting tools. Significant
advancement has been made in understanding toughening
mechanisms and some of these have been applied to improve the
toughness in the range 8-15 MPam1/2 From the view point of
multiphase ceramics, the flexural strength and fracture toughness of the
matrix materials can be enhanced by incorporating second phases 51
,52. The addition of hard secondary phases such as TiC, Ti(CN), WC
and SiC to alumina matrix provides great improvement in mechanical
properties 54, 55 , 61 , 62, 63, 64. The most important Al2O3 based
composites are those which contain TiC and ZrO2. Their properties
compared to pure alumina are shown in table 2.3.
Alumina-TiC composites
The hardness of alumina has also been shown to increase by adding
between 30 % and 40 % of TiC53. Such additions improve both hot
hardness and room temperature hardness but reduces the fracture
toughness. The increased hardness makes it more suitable for finishing
operations and for machining harder steels. The colour of this type of
Ceramic is black and is known on the market as Black Ceramic 53 .
K.F. Cai et al 54 investigated the effect of TiC additions on the
properties of alumina. Both the hardness and fracture toughness
A.T.Shonhiwa 22
Literature review
increased with increase in TiC content up to 30% vol and could
certainly, be extrapolated to still higher values. The increase in
hardness with increase in TiC could be explained by the fact that TiC is
relatively harder than Al2O3 and increase in fracture toughness was
attributed to effects of crack deflection and grain bridging by TiC grains.
X.Q You et al 55 investigated the effect of grain size on mechanical
properties and thermal shock resistance of Al2O3-TiC composites. In
addition to the remarkable improvement in mechanical properties
imparted by adding TiC into alumina composites they also found out
that decreasing grain size of Al2O3 and TiC results in improved thermal
shock resistance.
Alumina –ZrO2 composites(ZTA)
Zirconia toughened alumina (ZTA) is a well known two-phase binary
ceramic formed by adding ZrO2 powder to Al2O3 powder and sintering
to form a dense product with improved toughness over conventional
alumina ceramics 26. Thus the development of Zirconia-toughened-
alumina (ZTA) composites was aimed to substitute alumina ceramics in
applications where a higher fracture resistance is required 2. The
presence of second phase zirconia results in an enhancement of
flexural strength, and fracture toughness mainly attributed to the stress-
induced phase transformation. Work done previously has shown that
volumetric fraction of added zirconia is directly related to fracture
resistance of ZTA composites56, 57 ,58 59,. However it has also been
shown that addition of ZrO2 to an alumina matrix results in a hardness
decrease60. Such effect is associated with the lower hardness of
zirconia compared to that of alumina.
A.T.Shonhiwa 23
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Table 2.3 Al2O3 composites of commercial importance
Al2O3
30% TiC
Al2O3
40% TiC
Al2O3/ZrO2 Al2O3
Density (g/cm3) 4.22 4.36 4.01 3.9
Hardness( HV) 1810 1820 1700 2000
Bending strength
(MPa)
600-800 600-800 450 300-600
KIC( MPa m1/2
) 5.4 5.2 4.5 3-4
Λ, W/mk 30 35 15 20-30
α, 10-6
K-1 7 7 7 7-8
Rspec 103 Ωcm 4 2 1011 109
Application Cutting
tool
Cutting tool Cutting tool &
wear parts
Cutting tool &
wear parts
Alumina-WC composites
Lin Wang et al 61 investigated the influence of adding WC particles on
the Mechanical properties of alumina-matrix composites. In their work
they found out that dispersing tungsten carbide particles as a second
phase in an alumina matrix results in a composite with improved
mechanical properties. Flexural strength and fracture toughness of the
Al2O3 -WC composite (6 vol% WC) sintered at 1450 o C reached 581
MPa and 5.13MPam1/2 respectively 61. They concluded that a decrease
of the matrix grain size (due to pinning effect of WC grains) contributed
to increase in strength and presence of WC grains resulted in a more
tortuous crack-path which delays crack propagation leading to increase
in fracture toughness. Wilson et al 62 obtained fracture toughness of
7.1 MPam1/2 by hot pressing alumina and WC (20 vol%) at 1450 o C in
flowing argon. The hardness obtained in this work was ( 19MPa) which
is adequate for cutting tool applications, and is comparable to those of
A.T.Shonhiwa 24
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other hot –pressed materials. Microstructural investigations showed that
toughening was also due to crack deflection around the homogenously
distributed WC in the alumina matrix.
Alumina –(TiW)C composites
Wilson et al 63 also investigated the effect of mixed carbides as Ti and
W carbides on the mechanical properties of alumina matrix. Both
Titanium carbide and Tungsten carbide have been extensively studied
as reinforcing components in alumina ceramic matrix. Titanium carbide
and Tungsten carbide have room temperature hardness values in the
range of 18-23 and 17-22.5 GPa respectively 55, 63 .
Alumina-SiC composites
Young et al 64 fabricated Al2O3-SiC composites containing variable
amounts of SiC particles (5 to 30vol %) dispersed in an alumina matrix
and tested its potential as a cutting tool. The Al2O3 composites
exhibited higher hardness compared to the unreinforced matrix
whereas the fracture toughness remained practically constant up to
10% loading of SiC. Cutting tests done revealed that introduction of SiC
results in increased hardness and decreased grain size of the material,
thereby greatly improving its cutting performance, compared to the
commercial tools made of monolithic Al2O3 and Al2O3-TiC composites.
Reaction bonded aluminum oxide (RBAO)
Most aluminas sinter to full density at temperatures not lower than
1500 o C unless there are some additives. In addition there is some
shrinkage associated with high temperature sintering of alumina.
Reaction bonded aluminum oxide (RBAO) is a technique pioneered at
A.T.Shonhiwa 25
Literature review
Technische Universitat Hamburg-Harburg 65 to produce dense alumina
matrices at much lower temperatures with minimal shrinkage.
In this technique attrition milled Al/Al2O3 powder compacts are heat
treated in air such that all the Al gets oxidized into Al2O3 which then
sinters and bonds with the original Al2O3 to form a dense monolithic
alumina at temperatures much lower than that needed to conventionally
sinter alumina on its own. Another advantage is that oxidation of Al into
Al2O3 is accompanied by a volume increase which can compensate
firing shrinkage, thus enabling near net shaping. Other advantages of
this technique over conventional sintering of alumina include high green
strength, glass-phase-free grain boundaries and easy adaptability to
incorporation of second phase without causing the harmful residual
stresses normally encountered with shrinking matrix materials 65,67,68.
Work done by Claussen et al 66has shown that strength of reaction
bonded aluminum oxide bodies is comparable to that of conventionally
sintered dense alumina fired at higher temperatures.
Processing parameters affecting properties of RBAO bodies.
In an ideal reaction bonded aluminum oxide process all or most of the
Al should be converted into Al2O3 resulting in a homogenous matrix.
However in practice many problems are encountered e.g cracking,
bloating and incomplete reaction. Some of these problems can be
avoided if proper processing parameters are used. D. Holz et al 67
investigated the effect of processing parameters on phase and
microstructural evolution in RBAO ceramics. They observed that
important parameters controlling the reaction bonding of Al2O3 are Al
content, particle size and morphology of starting powders and green
density. In principle high Al content is necessary in order to exploit the
outstanding properties of RBAO (i.e the volume increase hence near-
net-shaping) Suvaci and Messing68 concluded that the maximum
A.T.Shonhiwa 26
Literature review
aluminum content of the precursor powder is limited to 60 vol % and for
aluminum contents above 60%, samples could not be completely
oxidized and would exhibit a dense oxide crust and an aluminum rich
core at the end of the reaction-bonding stage.
In general size of the Al particles affects strongly the microstructural
and compositional homogeneity of the sintered components 68, 69. Large
Al particles make fast and complete oxidation difficult and this can lead
to microstructural failures and heterogeneous particle distribution. The
critical aluminum particle size (i.e the largest aluminum particle size that
can be used to obtain dense ceramic materials via the RBAO process)
was determined to be approx 1.5μm 68.
E Suvaci and G. L. Messing 70 investigated the effect of initial α-Al2O3
particle size in RBAO process on the phase transformation of
aluminum- derived γ-Al2O3 to α-Al2O3. They concluded that γ-Al2O3 to α-
Al2O3 transformation have significant effect on subsequent sintering
kinetics, temperature and microstructure71, 72, . They demonstrated that
coarse α-Al2O3 used in normal RBAO process does not result in a
sufficiently high nucleation to affect the γ-Al2O3 to α-Al2O3 phase
transformation and seeding with fine α-Al2O3 results in high nucleation
frequency which have the resultant effect of reducing the
transformation temperature to as low as 963 o C. This smaller particle
size of the seeded RBAO Al2O3 decreases the sintering temperature to
1135 o C. Thus demonstrating that seeding is a viable method to tailor
the phase transformation and subsequently improve densification in
RBAO process.
The green strength of RBAO compacts can be attributed to the
presence of aluminum. On milling the Al particles are plastically
deformed, resulting in strong Al/Al contacts bridging the Al2O3 particles.
This results in high green density and green strength compared to that
of conventional ceramic green bodies73 . The pore system in RBAO
A.T.Shonhiwa 27
Literature review
compacts controls the oxygen flux and thus the oxidation reaction. This
means that although high green density of compacts favors low
shrinkage and near-net-shape forming it makes oxidation reaction
difficult and would result in incompletely reacted pellets.
Milling of RBAO powders
Since properties of reaction bonded aluminum oxide ceramics is
dependent on the size, morphology and homogeneity of the precursor
powder it means milling is the central processing step which determines
the properties of the powder and eventually the fired component. It
would be extremely harmful for the reaction bonding process if only a
fraction of the Al particles are effectively milled while the rest survive
because of undesired parameters or equipment setting. Important
parameters determining the quality of RBAO precursor powders are (a)
the volume ratio of the powder mixture to milling balls. (b) the rotation
speed (c) the milling time and (d) type of milling solvent 67.
F. Essl et al 69 investigated the effects on the milling efficiency of
milling medium and crystal size of the abrasive component. In their
work they compared organic solvents of different polarities namely,
ethanol, acetone and cyclohexane. They concluded that cyclohexane
(which is non polar) is more efficient than the two polar solvents ethanol
and acetone. They explained their finding in terms of stability. In the
non-polar cyclohexane there is no stabilization and so the slurry tends
to form powder agglomerates which are assumed to enhance the
milling efficiency.
They also assessed the effect of particle and crystal size on milling
efficiency. It was established that altering particle and crystal size does
not have an effect if milling is being done in non polar cyclohexane.
However in polar solvents milling efficiency is significantly lower with
fine alumina (median particle size: 0.5μm, crystallite size < 0.1μm) than
A.T.Shonhiwa 28
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when using coarser alumina (median particle size 18μm, size of
crystallites 1-2μm).
II-Soo Kim and Sang74 investigated the effect of different ball size and
alumina type on the processing of RBAO ceramics. The two types of
aluminas compared were fussed alumina and calcined alumina and
with all other parameters being constant they found out that fused
alumina is ground more effectively compared to calcined alumina and
this was attributed to the morphology of the former being sharp edged.
The investigation of different ball size distribution involved comparing
3mm balls with 50% 3mm balls + 50% 5mm balls. They observed that
small single size balls were more efficient compared to mixed size balls.
This was explained by considering the fact that probability of contact
between ball and powder is higher using small single size balls than
mixed size.
Mechanism of oxidation of Al in RBAO
In order to tailor the final properties of a RBAO body it is necessary to
understand the mechanism with which Al is being oxidized into Al2O3.
The RBAO process has been intensively characterized by dilatometry
and thermogravimetry 75. Typical TGA/DTA and dimensional changes
for the reaction bonded aluminum oxide process are shown in figure
2.4. The oxidation of Al starts at low temperatures ( >350 o C). Before
oxidation of Al heating of RBAO compacts is associated with mass loss
due to evaporation of organic phases and decomposition of hydrolysis
products, boehmite and diaspore. At temperatures below 450 o C the
reaction products are mainly amorphous Al2O3 and some traces of
crystalline γ-Al2O3 and above 450 o C Al oxidizes directly to crystalline
γ-Al2O3 and the preexisting amorphous phase also crystallizes to γ-
Al2O3. Maximum oxidation occurs at around 520 o C when there is
rapture of Al particles caused by decomposition of boehmite and
A.T.Shonhiwa 29
Literature review
diaspore. After this stage oxidation rate slows down again (between
520-660 o C ). Above 660 o C (melting point of Al) the rate increases
again. This can be explained by poor wetting of Al2O3 by liquid Al
resulting in liquid Al spilling into the void spaces where it gets readily
oxidized. This process continues until all Al is oxidized.
Figure 2.4 Typical TDA/TGA and dimensional changes for reaction
bonded aluminum oxide process75
The suggested reaction and mechanisms by Nils Claussen et al 73 have
been strongly supported by kinetic and thermodynamic studies75.
Based on isothermal reaction data, it was demonstrated that the
reaction kinetics in RBAO process follow a parabolic rate law and
A.T.Shonhiwa 30
Literature review
reaction rate depends strongly on the particle size of Al and is
controlled by oxygen diffusion. The activation energy of the process
before the melting point of Al (660 o C) was 112KJ/mol compared to
26kJ/mol above the melting temperature. This fact is evidence that the
reaction changes from gas-solid to gas –liquid after the melting of Al
Before melting of Al the gas is diffusing through along the porous grain
boundaries and after the melting point the molten Al spills into the void
space . In this situation there would be direct contact between oxygen
and Al resulting in overall reduction in activation energy of the process.
Solid State Chemistry for oxidation of Al in RBAO
In many gas-solid reactions the solid is porous, allowing diffusion and
reaction to take place simultaneously through out the solid. Thus the
reaction can be considered to take place at a diffuse zone rather than at
a sharp boundary as is the case in non-porous solids. Most of the work
done on solid state reactions has been based on the shrinking
unreacted core model 76
77.
Since many solid reactants have some initial porosity and the simple
shrinking unreacted core model is often inapplicable to such systems,
there have been efforts to find valid models for these reaction systems.
In general heterogeneous reactions involving a porous solid and a gas
generally include the following steps78.
1. Diffusion of the reactant gas within the pores of the solid.
2. Chemical reaction of the solid with the gas.
3. Diffusion of the gaseous product (if any) from the solid.
Thus as in any other chemical reaction system it is very important to
understand the relative significance of these steps in order to
understand the overall kinetics of the reaction system.
A.T.Shonhiwa 31
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Unlike in nonporous solids where there is a sharp boundary between
the unreacted core and the completely reacted layer in the case of
porous solids there is a gradual change in the degree of conversion
through out the particle. The external layer will be completely reacted
after a certain time and the thickness of this completely reacted layer
will increase toward the interior of the particle. Under these conditions
in contrast to non-porous solids, the reaction within the partially reacted
zone occurs simultaneously with diffusion of fluid reactants in this zone.
Thus the problem reduces to determine which of the processes
diffusion or reaction is rate determining.
If diffusion is rate determining, the reaction will occur in a narrow
boundary between the unreacted and completely reacted zones. If on
the other hand reaction is the rate determining step then the
concentration of the gas will be constant through out the solid and the
reaction will take place uniformly through out the solid 78, 79 .
S. P. Gaus et al 80 developed a model for the reaction bonded
aluminum oxide process which indicates that the process is controlled
by a combination of reaction and diffusion. In their model they plotted
the conversion profile of Al as a function of position within a pellet for
various reaction rates. In the case where the reaction rate was too low it
was seen that the concentration profile of Al was constant through out
the pellet. This is explained by the fact that at low reaction rates there is
enough gas through out the pellet and so the reaction proceeds
uniformly through out. Plot of aluminum with distance in this case was
as shown in figure 2.5a.
In the other extreme where the reaction rate was too high the reaction
occurred on the external layer leaving an unreacted core that shrinks as
the reaction proceeds. In this case concentration gradient of aluminum
was as shown in figure 2.5 b.
A.T.Shonhiwa 32
Literature review
Figure 2.5 a. Concentration profile
of Al oxidized as a function of distance
when reaction is rate determining80.
Modification of the RBAO process.
In order to reduce the shrinkage and even to achieve net-shaping the
RBAO process can be modified in various ways by incorporating other
metal or ceramic additives that exhibit a larger volume expansion on
oxidation. For instance Zr is associated with a volume expansion on
oxidation of 49%, Ti 76%, Cr 102%, and Nb 174% 73.
Claussen et al 73 showed that the microstructure and mechanical
properties of RBAO can be improved by incorporating 5-20% vol ZrO2.
The grain size of the matrix was shown to decrease with increase in
ZrO2 This is due to the fine distribution of ZrO2 particles at grain
boundaries. This hinders grain growth in the sintering stage.
Garcia et al 81 processed RBAO powders consisting of Al (40%),
Nb2O5-stabilised ZrO2 (Nb-OZP, 20%), Al2O3 (30 vol %) and Nb (10 vol
%) by the standard RBAO route. In spite of the high compression
pressures used (900 MPa) all the al and Nb oxidized below 900 o C. Nb
Figure 2.5 b. concentration profile of Al
oxidized as a function of distance when
diffusion is rate determining80.
A.T.Shonhiwa 33
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assisted in reducing the sintering shrinkage by a larger volume
expansion on oxidation. An interesting feature of this material was the
formation of needle-like grains consisting of Zr and Nb oxides. These
needle-like grains act as reinforcement particles by mechanism of crack
deflection and bridging resulting in improvement in fracture toughness.
S. Scheppokat et al 82 tested TiC and TiN as candidate materials for
particle reinforcement for RBAO. Since both TiC and TiN are not
sufficiently oxidation- resistant to withstand the heating cycle needed to
completely oxidize Al in RBAO process some modifications were done
so as to retain TiC and TiN after sintering.
In the case of TiC containing composite a precursor powder was
prepared which contained TiO2 to act as an oxygen donor for Al. The
presence of TiO2 eliminated the need for an oxidation step in air and
allowed the TiC pre-added to the composition to remain unoxidised.
The composites were fired in argon without a prior oxidation step. The
amount of TiO2 added was calculated so as to be enough to completely
oxidize the available Al.
For TiN containing composites advantage of an exchange reaction
between TiO2 and AlN to form Tin and Al2O3 was utilized. Thus in this
case powders containing TiN and AlN were heat treated in air up to 700
o C At this stage all the TiN was converted to TiO2. From this stage the
atmosphere was then changed to argon so that the formed TiO2 reacts
with the pre-added AlN to form Al2O3 and TiN. These samples achieved
a 96% density flexural strength of 280MPa and fracture toughness of
3.3 MPa m1/2
A.T.Shonhiwa 34
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2.3 Boron nitride
Boron nitride is a synthetic material which although discovered in the
early 19th century was not developed as a commercial material until the
latter half of the 20th century. Boron and nitrogen are neighbours of
carbon in the periodic table and their atomic radii are also similar to that
of carbon. It is not surprising therefore that boron nitride and carbon
exhibit similarity in their crystal structure.
In the same way as carbon exists as graphite and diamond, boron
nitride can also be synthesized in hexagonal and cubic forms.
Cubic boron nitride is the second known hardest material after
diamond. The compound crystallizes with a zinc blende structure which
closely resembles that of diamond with boron in 000 and nitrogen in ¼
¼ ¼ in a fcc lattice 95. Thus each atom is tetragonally linked to four
neighbouring boron or nitrogen atoms by strong covalent bonds as
shown in figure 2.6. It is this strong covalent bonding which is
responsible for the extreme hardness in both diamond and cBN, and it
also gives a reasonable explanation for the somewhat lower hardness
of cBN in comparison with diamond83.
Properties of cubic boron nitride.
Most of the properties of cubic boron nitride are similar to those of
diamond because of their electronic and structural properties which are
similar. Cubic boron nitride is the second hardest known material after
diamond. Compared to diamond cBN has several other advantages, in
particular higher thermal stability, stronger chemical stability with
respect to ferrous alloys, the possibility of n or p doping and the
emission of blue light at the p-n junction 84. The table below
summarizes the most important properties of c-BN and diamond.
A.T.Shonhiwa 35
Literature review
Figure 2.6 Crystal structure of cubic boron nitride85
Table 2.4 Physiochemical properties of cBN and diamond 84.
cBN Diamond
Structure Cubic F43m Cubic Fd3m
Unit cell parameter (Å) a=3.165 a=3.567
Interatomic distance (Å) d=1.57 d=1.54
Density (g cm -3) 3.48 3.52
Hardness( Kg mm -2) 4500 9000
Conductivity (W cm K) 13 20
Expansion (oC-1) 4.8 3.5
Stability against oxidation (oC) 1200 600
Graphitization (oC) >1500 1400
Resistivity (Ω cm ) 1010 1016
Refractive index (5893 Å) 2.117 2.417
A.T.Shonhiwa 36
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Synthesis of cubic boron nitride
Cubic boron nitride is normally synthesized from the hexagonal
polymorph, hBN by applying high pressures and high temperatures. In
principle it is possible to synthesize cBN from hBN by solid state phase
transformation at high pressure and high temperature without using any
catalysts86. However for practical purposes solvent catalysts are used
to reduce the high pressure high temperature conditions. The first
successful synthesis of cBN by high pressure and high temperature,
similar to that of diamond using Li3N as catalyst was first reported by
Went orf in 1957 87 . Since then various kinds of catalysts, which
reduce the temperature and pressure needed to synthesize cBN, have
been examined by various authors88 89.
Approximately fifty different kinds of catalysts such as alkali and
alkaline earth metal nitrides, fluoronitrides and ammonium borates have
been found to have catalytic effect on the formation of cBN. Most of
these materials form a eutectic with boron nitride and formation of cBN
proceeds via the dissolution of hBN in the eutectic liquid followed by
precipitation of cBN in its thermo dynamical stable region90 .Some of
the materials known as effective catalysts for cBN synthesis are
Lithium, magnesium, calcium, and their nitrides or boron nitrides (e.g
Li3BN2, Mg3BN3 and Ca3B2N4 ) 91 89 .
A.T.Shonhiwa 37
Literature review
Boron nitride Phase transformations.
Since boron nitride exists in two major crystalline forms, hexagonal
Boron Nitride and Cubic Boron nitride, intensive research has been
done to elucidate the temperature – pressure conditions under which
such transformations take place.
A phase diagram of boron nitride was first reported by Bundy and
Wentorf in 196392 , with cBN being the stable phase at ambient
conditions. This was changed however, in an ensuing publication by
Corrigan and Bundy in 197593, who considered a phase transition line
parallel to that of graphite/diamond, thereby intersecting at room
temperature at 1.3 GPa and thus implying now cBN at low pressure a
meta stable phase. This picture was adopted for many years until
calculations by Solozhenko 94 demonstrated cBN as the stable phase.
In order to further clarify the discrepancies on the cBN phase diagram
G. Will et al 95 also did some investigative in situ diffraction
experiments at high temperature and pressure to see the
transformation from hBN to cBN and the back transformation from cBN
to hBN. In their work they concluded that the phase diagram of boron
nitride is not comparable to that of carbon as was assumed in the past.
They also concluded that the cubic phase is definitely the stable phase
at low pressures and that the transformation depends strongly on
parameters like grain size, defect concentration and purity of the
starting material.
A.T.Shonhiwa 38
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Factors affecting transformation boron nitride.
H. Lorenz et al 96 investigated the influence of initial crystallinity on high
pressure –high temperature transformation of boron nitride by using
boron nitride powders with different initial crystallinity. They observed
that for the most disordered material (amorphous BN) the direct
transformation to cBN starts at relatively low pressures and
temperatures ( 1200 oC and 7.2 GPa ). This meant that decreasing BN
crystallinity of the initial material leads to a considerable reduction of the
thermodynamic conditions needed to produce cBN.. HRTEM
investigations showed that an amorphous layer forms around the
growing cBN grain and this enhances the diffusion of the growth units to
the interface of the new phase.
H. Sachdev et al 97 investigated the effects of grain size and impurities
on the cBN to hBN transition. They found out that powders with smaller
grain size transform into hBN at much lower temperatures compared to
powders with bigger grains. Three different powders with grain sizes of
0.75 -1.5μm, 40-80μm and 600μm had conversion temperatures of 900,
1300 and 1500 oC respectively. This can be explained by the fact that
powders with smaller grain size have a higher surface to- bulk ratio
therefore would react much quicker than larger crystals 97 . They also
found out that presence of impurity (boron oxide) on the surface of cBN
also influences the conversion mechanism. Conversion of cBN to hBN
proceeds by forming an intermediate rhombohedral phase and it is
considered that boron oxide acts as a catalyst for the formation of
rhombohedral boron nitride 97 .
A.T.Shonhiwa 39
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Oxidation of boron nitride
One advantage of cBN over diamond is its resistance to high
temperature oxidation. However at elevated temperatures boron nitride
would also start to oxidize. Thus in order to predict high temperature
properties of Boron nitride based composites it is necessary to have
an understanding of its behavior at elevated temperatures. V.A.
Lavrenko and A.F. Alexeev 98 did some work to investigate the high
temperature oxidation of various Boron nitride samples in the
temperature region of 600- 1200 oC. They found out that in the
temperature range 600 - 800 oC boron nitride does not react with
oxygen. Boron nitride oxidation was observed to start at 900 o C and
oxidation kinetics behaved in a parabolic manner until 1200 o C then it
behaved in a linear manner. Oxidation product was mainly B2O3
observed as thin film covering the original material and NO2 gas. N.
Jacobson and S. Farmer 99 also studied high temperature oxidation of
boron nitride in the temperature range 900- 1200 o C. They found out
that the main oxidation product was B2O3(l) and that the oxidation
kinetics are sensitive to crystallographic orientation , porosity and
impurity levels.
Boron nitride cutting tools
Cubic boron nitride has high hardness and thermal conductivity second
only to diamond and a low affinity to ferrous materials hence it is widely
used in grinding and cutting applications for ferrous materials. CBN is
normally used as a cutting material when hard metals become limited in
the cutting speeds that can be employed. This applies to hard work
piece materials such as high speed steel, tool steels, case hardened
steels, chilled cast iron, satellite etc. It offers no advantages and does
A.T.Shonhiwa 40
Literature review
not perform well on soft steels, inconel and nimonics and austenitic
stainless steel 53 .
PCBN works by self-induced hot cutting, a process which causes local
softening in the shear zone of the work piece material. The heat
generated is discharged through the chips as well as through the PcBN
insert, leaving the hardness of the work piece unaffected. In achieving
effective self-induced softening of the work piece, PcBN‘s high hot
hardness and chemical stability are fully exploited. Good toughness and
cutting edge stability are also important under these machining
conditions 100.
Typical cBN Cutting tools
Conventional cBN based composites available commercially as cutting
materials contain a ceramic or metallic binder which facilitates sintering
and optimize cutting performance. Usually metals of group 4, 5 and 6 of
the periodic table and/or other metallic elements, such as aluminum,
cobalt and nickel are used to activate sintering101, 102 . Table 2.5 shows
properties of typical cBN cutting tools with different binders.
CBN cutting tools can also be available as solid indexable inserts or as
inserts consisting of an upper face of cBN laid onto a hard metal base.
It can also be available as a piece of cBN brazed onto a corner of a
hard metal indexable insert. Solid indexable inserts are particularly
suitable for heavier roughing work, especially for machining tools 53.
A.T.Shonhiwa 41
Literature review
Table 2.5 Typical properties of cBN cutting tools103
Wt %
cBN
Binder
phase
Elastic
modulus(GPa)
Poisson’s
ratio
cBN grain size
(μm)
90 AlN and
AlB2
648 0.145 15
80 TiC and Al
compounds
981 0.160 10
50 TiC and Al
compounds
595 0.170 1-2
45 TiC and Al
compounds
582 0.177 1-2
PCBN Physical and Mechanical properties.
The physical properties of PCBN cutting tool materials is determined by
a number factors. The cBN grain size and content can affect the
abrasion resistance, edge quality and thermal conductivity of the PCBN
product. For machining conditions where the mode of wear is
predominantly abrasion , for example in the machining of cast iron
coarse grain, high cBN content material will give optimum performance
and for finish machining where good edge quality is required, fine
grained lower cBN content material would be appropriate 100. Type of
binder phase used also has a considerable effect on the performance of
PcBN materials, particularly in the case of low cBN content products
where the PCBN binder content can be less than 50% by volume.
Figure 2.6 shows the relationship between cBN content and grain size
on PCBN properties of some commercial PCBN cutting tools.
A.T.Shonhiwa 42
Literature review
Figure 2.7 Effect of cBN content and grain size on PCBN properties104
Recent developments in cBN cutting tools
Benko et al 101 studied the phases and microstructures developed in
cBN composite synthesized by hot pressing with Al as the binding
phase in the molar ratio BN;Al of 9;1. They observed that BN and Al
reacted to form AlN and that in the area between AlN and BN phases
polycrystalline AlB10 and AlB12 phases could be seen. Hardness of the
samples increased after annealing and thermal treatment also resulted
in increase in mechanical strength of the sintered BN-Al system.
In another investigation Benko et al 105 also looked at the phases and
properties of composites cBN-TiN and cBN-TiC. Experimental cBN-
TiC/TiN composites were prepared by high pressure hot pressing and
samples were subsequently heat treated It was shown that in the
45
50
90
80
1μm 2μm 6μm 8μm
cBN average grain size (microns)
Improved edge quality
cBN
vol%
Abrasion resistance
Thermal conductivity
DBN45 DBN50 DBA80
DBA90
A.T.Shonhiwa 43
Literature review
temperature range 1000 to 1400 oC TiN reacts with cBN to form one
new phase , TiB2 and that TiC reacts with cBN forming two new phases
TiB2 and TiC0.8N0.2.
On the other hand, a polycrystalline cBN compact with no additives
(PCBN) can also be synthesized by direct transformation from a low
pressure phase of BN under ultra high pressure and high
temperature106 ,107 T. Ohashi et al 102 studied the influences of
synthesizing conditions on properties of PCBN such as microstructure,
hardness, cutting performance and thermal conductivity. In their work
hBN was used as starting material and was converted to cBN by
treating it at 6.8 GPa and at temperatures from 1800 o C to 2500 o C.
Samples treated at 1800 o C consisted of hBN and cBN started to form
at 2100 o C. At 2100 o C the microstructure consisted of fine (<1μm)
homogenous cBN grains and became heterogenous with some coarse
grains at 2300 o C and at 2500 o C there was remarkable grain growth
with twinning. It was observed that hardness decreased with increase
in processing temperature, probably due to grain growth. Thermal
conductivity increased with increase in processing temperature,
probably due to the fact surface area decreases with increase in
processing temperature. For cutting tests results it was found out that
sample processed at 2100 o C had excellent wear resistance while
sample processed at 2500 oC had similar wear resistance to
conventional cBN-Co composites.
A.T.Shonhiwa 44
Experimental Details
.Chapter 3: Experimental Details
This chapter give details of the raw materials used, equipment and
analytical techniques employed for characterizing materials in this
thesis.
3.1 Chemicals
Chemicals used in this study and their respective specifications are
shown in table 3.1.
Table 3.1 Chemicals used.
Chemical Supplier Grade and properties
Aluminum Saarchem (R.S.A) 99.8% purity with particle size of 2-5microns
Alumina Sumitomo (Japan) Α-Alumina (AKP50) with a purity of 99.99%
and particle size of 0.1 -0.3 microns
Cubic Boron nitride Element Six (RSA) Particle size 2 -5 microns
Cyclohexane Associated Chemicals (RSA) Assay purity of 98.5%
Stearic acid Hopkin and Williams (UK) Iodine value < 4% and an acid value of 200 to
100. Sulphated ash <0.1%
Zinc stearite Riedel-de-Haen (Germany) Assay of Zinc 10-12% and ash content of 12-15%
A.T.Shonhiwa 45
Experimental Details
3.2 Equipment
Attrition mill
A Szegvari attritor system type B from Union Process was used for
milling powder raw materials. The attritor was water cooled fitted with a
750cc alumina vessel and an alumina agitator. In this work 2mm
alumina balls were used as milling media and cyclohehane was used
as milling solvent.
Rot vapor
A Heidoiph Laborota 4000 rot vapor fitted with a hot water bath was
used to dry off the solvent (cyclohexane) from the milled slurry to form
powder.
Uniaxial press
A custom made uniaxial hydraullic press with a 30mm radius plunger
capable of delivering a pressure of 40MPa was used to make the
pellets. Although the press plunger is capable of delivering a maximum
pressure of 40MPa higher pressures were achieved by using dies with
smaller cross sectional area. For this work a stainless die with an
internal diameter of 18mm was used .
Box furnace
A Eurotherm 2416 box furnace manufactured by Elite Thermal systems
capable of heating up to 1600 o C was used to heat treat the pellets so
as to convert aluminum into alumina. To avoid any possible reaction or
cross contamination all samples to be fired were placed in alumina
boats.
A.T.Shonhiwa 46
Experimental Details
Hot press
A unaxial hot press system was used to densify the samples. This
system is equipped with a carbon heated furnace with a maximum
obtainable temperature of 2000 oC and a steel loading frame capable
of delivering loads of up to 10 000 Kg. This force is applied unaxially
through graphite punches. In addition the system is also equipped with
a water cooling system and a vacuum pump capable of attaining
vacuum levels of 10 mtorr. All samples were either heated in Argon
atmosphere or under vacuum.
Figure 3.1 Hot press system.
FURNACE
REACTION RAM
PRESSING RAM
CONTROL PANEL
INSTRUMENT PANEL
HYDRAULIC PUMP
A.T.Shonhiwa 47
Experimental Details
3.3 Analytical techniques
Phase analysis (Qualitative and quantitative) X-ray diffraction
Qualitative and quantitative phase analysis of powders, reacted pellets
and sintered samples was done using X-ray diffraction. For green and
reacted samples (partially sintered) measurements were done on the
surface and for sintered samples measurements were done on polished
cross section. Diffraction patterns were collected using a Bruker AXS
D8 machine equipped with a primary beam Göbel mirror, a radial soller
slit, a V Åntec-1 detector and using Cu-K α radiation (40kV, 40mA).
Data were collected in the 2θ range 5 to 90 o in 0.021 steps, using a
standard scan speed with an equivalent counting time of 14.7 s per
step.
For qualitative analysis the resulting diffractograms were analysed
using X’Pert high score software developed by Philips.
Quantitative phase analysis was done using the following equation
pmhkl KKI ) ………………………………………………………………(1)
Where I(hkl) is the peak intensity , Km represents the physical constant
and measurement factors (which is constant for all samples measured
on the same machine under the same conditions) and Kp represent the
phase related factors.
Kp for each phase was then calculated using the equation
hklhkl
hklhkl
hklF
V
pKp
sin2sin2
2cos1 2
2
2 ………………………………….(2)
A.T.Shonhiwa 48
Experimental Details
Where p(hkl) is the multiplicity, V is the cell volume and F is the
structure factor. The intensity I(hkl) for each peak was determined by
taking the peak area of the respective signal using a mathematical
software.
Having evaluated the intensity and Kp the volume fraction of each
respective phase was then calculated using the relationship
100*%
924
924
32
2
2
32
3
3
)(
OBAl
OBAl
BN
BN
Al
lA
p
OlA
OAl
hkl
Kp
I
Kp
I
Kp
I
K
I
Kp
I
V
OAl
OAl…………..(3)
Peaks and relevant parameters used for quantifying each phase are
shown in table 3.2.
Table 3.2 Phase parameters used for quantification.
Phase 2θ h k l P F2 V Kp Ref
Al2O3 35.14 104 6 6496 255 2.874 1-071-1123
Al 38.47 111 8 1283 66.9 9.025 1-089-2769
cBN 43.26 111 12 267.5 47.4 4.326 1-079-062
Al18B4O33 16.46 110 4 607.14 164 2.104 0-029-0009
Particle size determination -Malvern analyzer
A Malvern Mastersizer 2000 was used for particle size determination. A
small quantity of powder was diluted in ethanol in a small beaker, then
placed in an ultrasonic bath to break up any agglomerates. The
analyzer was filled with ethanol and then the diluted powder was slowly
added The accuracy of the equipment is ±4% (volume median
A.T.Shonhiwa 49
Experimental Details
diameter). In all instances the pump speed was set at 40% of its
maximum and the ultrasonics adjusted to 80% and a 45mm focal
length lens was used.
Microstructural analysis-Scanning electron microscope
A Joel scanning electron microscope was used to assess the
microstructure of the materials after reaction and hot pressing. The
microscope was normally operated at an accelerating voltage of 20 kV.
Before analysis samples were coated with gold to provide conductivity
and were the mounted on graphite tape before analysis.
Pore size and distribution- Mercury Porosimeter
Median pore diameter of green and reacted samples were determined
using mercury porosimetry (Poresizer 9320, Micrometrics, USA).
Density measurements.
Density of green and reacted samples was determined geometrically,
by measuring the mass, diameter and height of the pellets.
Density of sintered samples was determined using the Archimedes
method. Samples were weighed dry (Md) before being boiled in water
for three hours in order to drive air from the pores. After boiling the
samples were left soaking in the water overnight. The mass of the
samples suspended in water (Ms) was determined followed by the
soaked mass (Mw). Please note that before determining the soaked
mass the samples were first wiped with a light towel to remove excess
water from the surface of the samples. The density and porosity of the
samples was then calculated using the following equations.
A.T.Shonhiwa 50
Experimental Details
sw
d
MM
MDensity ………………………………………...(1)
100*dw
sw
MM
MMPorosity …………………………………..(2)
Where ever possible density values were expressed as a percentage of
the expected theoretical density. The expected theoretical densities
were calculated from the percentage mass composition of the
constituent phases and their respective theoretical densities using the
formular.
ρ
z
z
y
y
x
x
P
M
p
M
p
M %%%
100 ………………………………(3)
Where Mx, My and Mz are the percentage mass compositions of phases
x, y and z and ρx, ρy and ρz are the respective theoretical densities.
Hardness and Fracture toughness measurements
Hardness measurements were done using a Leco V-100 hardness
tester. All measurements were done using a load of 10 Kg for 10 Sec.
Hardness was calculated using the formular
22
4.1854a
PH v
Where
P = Applied load in newtons and 2a = Indentation diagonal in μm.
Fracture toughness determinations were done by measuring crack
lengths from the indented sample according to the equation given
below.
A.T.Shonhiwa 51
Experimental Details
l
PHK v
IC*4
**0889.0 (103)
Hv = Vickers hardness
P = Load and
l= c – a Where 2a is the indentation diagonal and 2c is the combined
length of two opposite cracks including the indentation diagonal length.
A.T.Shonhiwa 52
Experimental Details
3.4 Experimental procedure
Milling
Weighed quantities of alumina, aluminum and cubic boron nitride were
slowly charged into a 750 ml alumina vessel containing cyclohexane,
1mm alumina balls and 0.5% stearic acid. The amount of stearic acid
added was determined from the total mass of alumina, aluminum and
cBN .The vessel was attached to a Szegvari attritor system fitted
Al2O3 blades rotating at a speed of 50rpm. When all the powder has
been added the rotation speed was increased to 700rpm and the
powders were milled for eight hours.
In each instance the masses of alumina, aluminum and cubic boron
nitride charged were determined in such away as to give respective
volume percentages required for that particular batch. Densities of
alumina, aluminum and cubic boron nitride used for this determination
are 3.98, 2.7 and 3.48 g/cm3 respectively. . In each instance the
volume ratio of powder to balls was 1:5 and the vessel was filled up to
two thirds of its volume capacity.
The starting compositions (volume) of powder mixtures used in this
work are given in table 3.2 .
Please note that elsewhere in this work samples will be designated
according to their initial Al:Al2O3:BN compositions (volume %) before
milling. For an example 502030 refers to sample with the following
initial compositions of Al (50 volume %), Al2O3 (20 volume %) and
BN (30 volume %) and 505000 refers to sample with an initial
composition of Al (50 volume %) and Al2O3 ( 50 volume %).
A.T.Shonhiwa 53
Experimental Details
Table 3.3 Mass composition of starting materials
Compositions (Mass %) and designations
Material
502030 504010 454510 505000
Al
42.32 41.03 26.23 40.42
Al2O3
24.95 48.39 53.40 59.58
BN
32.73 10.58 10.38 0
Drying and sieving
After milling the slurry was transferred to a rotary evaporator with a
water bath set at ± 60 o C. The alumina balls used for milling were
thoroughly washed with cyclohexane and the resultant elutants was
also transferred to the rotary evaporator to maximize on the recovery of
milled powders . After drying the powders were passed through 38μm
sieve to break soft agglomerates.
Cold pressing.
The powders were then weighed out into ± 2.00gram quantities for
green compaction. Green compacts were produced using stainless
steel dies with an internal diameter of 18 mm and a custom made
uniaxial press described in section 2.2.3. For each powder composition
samples were pressed at three different pressures that is 45, 90 and
180 MPa.
A.T.Shonhiwa 54
Experimental Details
Heat treatment.(Oxidation)
The green pellets were reacted in air in a box furnace (Lenton)
described in section 2.2.4 in order to facilitate conversion of Aluminum
metal into the oxide
Al + 3/2 O2 Al2O3 .
In a typical heat treatment cycle the samples would be heated from
ambient temperature up to 500 o C at a rate of 3 oC per minute followed
by a soak at 500 o C for 300 minutes. From 500 o C the temperature
would increase at a rate of 3 oC to the required reaction temperature .,
Once the reaction temperature has been attained the furnace would
start to cool down to ambient temperature at rate of 10 oC per minute.
However for samples heat treated to 1000 oC a soaking time variable
was also introduced into the study whereby instead of stopping the
reaction after reaching 1000 oC samples were soaked for different
times i.e 0 , 60, 180 and 300 minutes to assess the effect of soaking
time on the reaction before cooling to ambient temperature.
After heat treatment samples were assessed for the following density ,
mass, dimensions (diameter and height), and phase analysis ( both
qualitative and quantitative) using X-ray diffraction as described in
section 2.3.2.
Hot pressing
A hot press equipment described in section 3.2.5 was used to densify
the samples. Reacted pellets were placed in 18mm diameter graphite
pots equipped with pistons of same diameter. Before sintering all
graphite components, including punch and die were coated with an hBN
suspension.
After placing the samples in the furnace the whole system was
evacuated to pressures less than 100 mtorr using a vacuum pump. This
A.T.Shonhiwa 55
Experimental Details
was followed by purging argon into the system at a rate of 40ml/min.
Once there was a stable flow of argon in the system the furnace was
switched on and rumped at a rate of 40 oC per minute to a temperature
which is 100 oC less than the intended sintering temperature. At this
point the hydraulic system was then activated and load was slowly
applied until the desired load was attained. Once the desired load was
attained the temperature was rumped to the desired value and soaked
for the required time. During the soaking period pressure was
maintained manually. After the soak period the load was removed and
furnace rumped down to ambient temperature at a rate of 10o C per
minute. Below is a schematic representation of a typical temperature –
pressure profile used.
Figure 3.2 Schematic representation of the heat treatment - pressing cycle used to densify samples.
120 mins
Soak
40 o C/min
10o C/min
2000Kg
1000Kg
1500Kg
500Kg
Load/Kg
1200 o C
600 o C
1300 o C
Temp/o
C
Time/mins
A.T.Shonhiwa 56
Results and Discussions
Chapter 4: Results and discussions.
Introduction
This chapter deals with the experimental results obtained at various
stages from milling of the raw materials through compaction into green
bodies and initial heat treatment up to final sintering. At each of these
processing stages various techniques are used to characterize
properties of the materials. These are discussed together with their
effects on the properties of the final product. Also discussed in this
chapter is the oxidation kinetics of Al into Al2O3 and the interaction of
cBN with Al2O3 together with the relevant phase diagrams and
thermodynamic considerations.
4.1 The milling Process
Wet milling of Al2O3 and Al raw materials is a very important step which
predetermines the characteristics of the precursor powders hence
properties of the final product. Success of the reaction bonded
aluminum oxide process depends to a large extent on the particle size
of the aluminum. Coarse aluminum particles result in incomplete
oxidation leading to microstructural in homogeneities. Thus the primary
objective of the milling process is to produce fine aluminum crystallites
and to have them homogenously distributed with the alumina and cBN
particles. Only Al and Al2O3 mixtures were attrition milled and after
milling and drying the right amount of cBN was then added to the milled
powder and sonicated in cyclohexane for ten minutes.
Predicting characteristics of the milled powders became a complicated
issue because the chemistry of one of the components (aluminum)
changed during milling i.e oxidation and hydroxylation.
A.T.Shonhiwa 57
Results and Discussions
Optimization of the milling process involved monitoring particle size
distribution morphology and surface area of the milled powders as a
function of milling time.
Particle size distribution.
Size distribution of the raw materials in the received state was
measured using Malvern Analyzer (see section 3.3.2 for details). The
powders were milled as described in section 3.2.1 with samples being
withdrawn periodically to check size distribution as a function of time
(see Figure 4.2).
Figure 4.1 Particle size distribution of aluminum and alumina before
Figure 4.33 Phase compositions for materials heat treated to 1000 oC in air followed by sintering to 1300 oC in Argon.
A.T.Shonhiwa 98
Results and Discussions
22 24 26 28 30 32 34 36 38 40 42 44 46 48 50
Al2O
3 + cBN
Al
Al2O
3
Al2O
3
Al2O
3
505000
454510
504010
502030
Position (o 2 Theta)
Figure 4.34 Phase compositions for materials heat treated to 800 oC in
air followed by sintering to 1300 oC in Argon.
A.T.Shonhiwa 99
Results and Discussions
20 22 24 26 28 30 32 34 36 38 40 42 44 46 48 50
Al
Al2O
3
Al2O
3
Al2O
3
Al2O
3 + cBN
505000
454510
504010
502030
Position (o 2 Theta)
Figure 4.35 Phase compositions for materials heat treated to 800 oC in air followed by sintering to 1300 oC in Vacuum.
A.T.Shonhiwa 100
Results and Discussions
Sintered densities.
Densities of sintered materials were determined by means of the
Archimedes method described in section 3.3.5 of this thesis. The
theoretical densities were calculated using the rule of mixtures, using
the following theoretical densities for the constituents phases Al 2.7,
Al2O3 3.98, cBN 3.48 and Al18B4O33 2.68g/cm3 . The volume proportions
of the various phases were determined using the quantitative X ray
diffraction method described in section 3.3.1 of this thesis. Densities of
sintered samples are shown in table 4.4.
Table 4.4 Density of sintered materials.
Sample Density g/cm3 Density (%)
502030 1000 1300 Ar 3.55 94.12
504010 1000 1300 Ar 3.70 95.45
454510 1000 1300 Ar 3.69 95.33
505000 1000 1300 Ar 3.86 97.38
502030 800 1300 Ar 3.65 95.67
504010 800 1300 Ar 3.77 96.40
454510 800 1300 Ar 3.78 96.57
505000 800 1300 Ar 3.85 97.22
502030 800 1300 v 3.68 96.50
504010 800 1300 v 3.79 97.02
454510 800 1300 v 3.79 97.19
505000 800 1300 v 3.95 99.67
A.T.Shonhiwa 101
Results and Discussions
Effect of composition on sintered density
For samples heat treated under the same conditions density decreases
with increase in cBN content. This can be explained by the hardness of
cBN and its resistance to plastic deformation which results in it not
participating in the sintering process. Figure 4.36 shows the effect of
composition on density for samples sintered under different conditions.
Figure 4.36 Effect of composition and heat treatment on density of sintered samples.
92
93
94
95
96
97
98
99
100
1000 1300Ar 800 1300Ar 800 1300V
Heat treatment
Den
sit
y (
%) 502030
504010
454510
505000
A.T.Shonhiwa 102
Results and Discussions
Effect of heat treatment on sintered density
For all materials density changes with heat treatment. Samples with
cBN, exhibit a relatively lower degree of densification if initially heat
treated to 1000 oC compared to when they are when initially heat
treated to 800 oC.
This lowering of density when samples reacted at 1000 oC can be
attributed to the presence of Al18B4O33 which has a lower density
(2.68g/cm3) compared to Al2O3 (3.98g/cm3). This slight decrease in
density is not realized in sample without cBN (505000).
Effect of sintering atmosphere on sintered density.
For all samples a major improvement in density is realized if samples
are sintered under vacuum compared to Argon,. This is because
towards the final stages of sintering some argon gas is trapped in the
pores and closure of these pores depends on the ease with which the
trapped Argon can diffuse into the surrounding matrix. A comparison of
densities obtained by sintering in Argon versus sintering under vacuum
is shown in figure 4.37. From figure 4.37 it can be seen that the density
obtained by sintering under vacuum is more pronounced in the sample
which does not have cBN (505000). This sample has the highest
density of 99.67%
A.T.Shonhiwa 103
Results and Discussions
Microstructural analysis
a. Alumina matrix.
Sample 505000 when sintered at 1300 oC consists of a homogenous
matrix with well defined grains in the submicron region. Figure 4.38
and 4.39 show micrograph of alumina matrices sintered in Argon and
vacuum at 1300 oC followed by thermal etching at 1200 oC in air.
Sintering in vacuum does not have any significant effect on
microstructure although it results in improved densification.
A.T.Shonhiwa 104
Results and Discussions
Figure 4.37 SEM image of sample 505000 reacted to 800 oC in air and sintered at 1300 oC in Argon.
Figure 4.38 SEM image of sample 505000 reacted to 800 oC in air and sintered at 1300 oC under vacuum.
A.T.Shonhiwa 105
Results and Discussions
b. Samples with cBN
Samples with cBN (502030, 504010 and 4545100 consist of cBN
particles (dark colour) evenly distributed within an alumina matrix
(lighter colour) as shown in figures 4.40 and 4.41.
In figure 4.42 (higher magnification) shows that there is no reaction
between cBN and surrounding alumina. This results in pores being
formed on the alumina cBN interface and this could be one of the
reasons why density decreases with increase in cBN.
Figure 4.39 SEM image of sample 502030 reacted to 800 oC in air and sintered at 1300 oC in Argon.
A.T.Shonhiwa 106
Results and Discussions
Figure 4.40 SEM image of sample 504010 reacted to 800 oC in air and sintered at 1300 oC in Argon.
Figure 4.41 SEM image of sample 454510 reacted to 800 oC in air and sintered at 1300 oC in Argon.
cBN
Pores
Alumina
A.T.Shonhiwa 107
Results and Discussions
Mechanical properties of sintered materials.
Hardness and fracture toughness of the sintered materials were
determined using the indentation method as described in section 3.3.6
of this thesis. Figure 4.43 is a typical indentation on a polished section
of one of the materials.
Figure 4.42 A typical indent produced on a 504010 sample using a 10 kg load for 10 seconds.
A.T.Shonhiwa 108
Results and Discussions
Table 4.5 Hardness and fracture toughness values for samples heat treated under various conditions
Sample Hv10 (GPa) KIC (MPa m1/2)
502030 1000 1300 Ar 22.2 ±1.2 3.9±0.5
504010 1000 1300 Ar 19.9±1.7 3.5±0.7
454510 1000 1300 Ar 20.1±1.1 3.3±0.8
505000 1000 1300 Ar 18.5±2.2 NA
502030 800 1300 Ar 24.5±1.6 3.9±0.8
504010 800 1300 Ar 21.8±0.9 3.3±0.9
454510 800 1300 Ar 20.7±0.8 2.8±1.0
505000 800 1300 Ar 18.5±1.1 2.0±0.9
502030 800 1300 v 24.6±0.9 3.9±1.2
504010 800 1300 v 22.1±1.1 3.2±1.1
454510 800 1300 v 21.8±0.8 3.2±1.0
505000 800 1300 v 20.2±1.2 2.8±0.7
Effect of composition on hardness
For all sample compositions hardness is seen to increase with increase
in cBN content. This is expected since boron nitride has a higher
hardness compared to alumina and hence by the rule of mixtures the
resultant hardness of a composite should increase with increase in
volume fraction of the harder phase. For an example for samples heat
treated at 1000 oC in air followed by sintering at 1300 oC in Argon
hardness increases from 18.58 GPa for material without cBN (505000)
to 22.24 GPa for sample with 30 % volume cBN (502030). Figure 4.44
shows the effect of cBN content on hardness for samples heat treated
under various conditions.
A.T.Shonhiwa 109
Results and Discussions
18
21
24
27
1000 1300Ar 800 1300Ar 800 1300V
Heat treatment
Hard
ness H
v 505000
454510
504010
502030
Figure 4.43 Effect of composition on hardness for samples heat treated under various conditions.
Effect of heat treatment on hardness
Hardness is also dependent on heat treatment . Samples heat treated
in air at 1000 oC followed by sintering at 1300 oC in Argon have slightly
lower hardness values compared to same materials which have been
heat treated at 800 oC in air followed by sintering at 1300 oC in argon.
For an example sample 502030 has a hardness of 24.5 GPa when heat
treated at 800 oC followed by sintering to 1300 oC compared to 22.2
GPa when initially heat treated to 1000 oC followed by sintering to 1300
oC. This compromise in hardness for samples heat treated at 1000 oC
can be attributed to the presence of Al14B4O33 phase.
Hardness values for samples heat treated in air at 800 oC followed by
sintering at 1300 oC in vacuum are higher than hardness of same
A.T.Shonhiwa 110
Results and Discussions
materials heat treated in air at 800 oC followed by sintering at 1300 oC
in Argon. Improvement of hardness for samples sintered under vacuum
compared to their counterparts sintered in argon can be ascribed to
increased densities achieved by vacuum sintering. Figure 4.45 shows
the effect of various heat treatments and composition on hardness.
15
18
21
24
27
0 10 20 30
cBN (% vol)
Hard
ness H
v
1000 1300Ar
800 1300Ar
800 1300V
Figure 4.44 Comparison of hardness values for samples heat treated under various conditions.
Effect of composition on fracture toughness.
Presence of cBN in an alumina matrix has resulted in improved fracture
toughness values. For an example for sample reacted at 800 oC in air
followed by sintering to 1300 oC in Argon fracture toughness has
increased from 2.0 MPa m1/2 for sample without cBN to 3.9 MPa m1/2 for
sample with 30 % volume cBN. Figure 4.46 shows the effect of cBN
content on fracture toughness for samples heat treated at 800 oC
followed by sintering at 1300 oC in Argon. This enhancement of fracture
A.T.Shonhiwa 111
Results and Discussions
toughness can be attributed to crack deflection by cBN particles as
shown in figure 4.47. As the cBN content increases it means the
propagating crack will follow a more tortuous path resulting in even
higher fracture toughness values.
0
1
2
3
4
5
6
0 10 20 30 40
cBN (% vol)
Fra
ctu
re T
ou
gh
ness(M
Pa m
1/2
)
Figure 4.45 Effect of cBN content on fracture toughness for samples
heat treated at 800 oC in air followed by sintering at 1300 oC in Argon.
Figure 4.46 Crack deflection around cBN particles in sample 502030 sintered at 1300 oC in Argon.
Crack deflection
A.T.Shonhiwa 112
Results and Discussions
Effect of heat treatment on fracture toughness
Considering same materials sintered under different conditions there
does not seem to be any correlation between fracture toughness and
heat treatment. For an example sample 502030 when heat treated to
800 oC in air followed by sintering at 1300 oC in Argon has a fracture
toughness of 3.96 MPa m1/2 ( highest recorded for all samples ) which
goes down to 3.90 MPa m1/2 when the sintering is done in vacuum.
Thus unlike density and hardness which increase on changing
atmosphere from Argon to vacuum , fracture toughness does not
improve with change of sintering atmosphere.
A.T.Shonhiwa 113
Conclusions and recommendations
Chapter 5
Conclusions and recommendations
5.1 Summary
This work has demonstrated that incorporating cubic boron nitride in a
reaction bonded aluminum oxide matrix results in composite materials
with improved hardness and fracture toughness compared to pure
aluminum oxide.
Reaction bonded aluminum oxide matrix.
Reaction bonded aluminum oxide (RBAO) was used instead of
conventionally sintering alumina. This required pre heat treating the
composites in an oxidizing atmosphere to facilitate oxidation of
aluminum into alumina without oxidizing cubic boron nitride into B2O3.
This required a thorough understanding of the oxidation kinetics of
aluminum. In particular the effects of the following factors on the
oxidation of aluminum in RBAO were investigated,
1. Compaction pressure
2. Temperature
3. Chemical composition
The effects of each factor on the oxidation of aluminum are
summarized below.
Compaction pressure.
Compacts were made at different pressures 45, 90 and 180 MPa. It
was observed that degrees of reaction (conversion of Al into Al2O3) did
not vary much with change in compaction pressure.
A.T.Shonhiwa 114
Conclusions and recommendations
Temperature.
For all sample compositions degree of reaction increased with
increase in temperature with maximum oxidation taking place in the
range 600 – 800 degrees.
Chemical composition.
Comparing materials with different cubic boron nitride loading
revealed that presence of cubic boron nitride inhibits oxidation of
aluminum. This can be attributed to the possibility of BN forming a
thin film of B2O3 in the vicinity of Al which then hinders diffusion of
air resulting in lower degrees of oxidation.
The oxidation investigations showed that heat treating samples
containing cubic boron nitride to higher temperatures (1000 oC)
results in the formation of B2O3 which then reacts with Al2O3 to form
Al18B4O33. Thus to avoid formation of B2O3 and Al18B4O33 it was
decided that the optimum temperature for oxidizing aluminum in
reaction bonded aluminum oxide process ( RBAO) is 800 oC.
Sintered samples
Samples reacted to 800 oC and 1000 oC in air were sintered to 1300
oC in vacuum and Argon for further densification.
It was observed that density decreased with increase in cubic boron
nitride. This could be related to the hardness of cubic boron nitride
and its resistance to plastic deformation.
For all samples there was an appreciable improvement in
densification if sintering was done under vacuum compared to
Argon.
For all samples, introduction of cubic boron nitride resulted in
appreciable improvement in hardness and fracture toughness. Both
A.T.Shonhiwa 115
Conclusions and recommendations
hardness and fracture toughness increased with increase in cubic
boron nitride loading.
Thus this investigation demonstrated that introducing cubic boron
nitride (up to 30% by volume) in a reaction bonded aluminum oxide
matrix results in a composite material which has enhanced
properties compared to pure alumina. This composite has
hardness and fracture toughness values of 24.6 GPa and 3.9 MPa
m1/2 respectively making it possible candidate for wear applications.
5.2 Future work
Sample 502030 when sintered for 2 hours at 1300 oC under vacuum
had the best mechanical properties ( hardness and fracture
toughness values of 24.6 GPa and 3.9 MPa m1/2 respectively) and
a density of 96.50%. Mechanical properties can further be enhanced
by increasing cubic boron nitride content to say 40%. This however,
is most likely to further reduce the density. Density can be improved
by increasing sintering temperature say to 1350 oC and sintering
time to 5 hours.
In order to fully understand the effect of cubic boron nitride on the
properties of alumina it might also be worth doing a detailed
microstructural characterization of the cBN/alumina interface region.
In particular high resolution electron microscopy (HREM) imaging
can provide considerable detail of the interface structure which can
then be used to explain enhancement in fracture toughness and
hardness.
A.T.Shonhiwa 116
Appendix
Appendix
Table A.1. Composition by volume of sample 502030 as a function of temperature and pressure.
PRESSURE
MPa TEMP
oC
TIME MINS
% Composition Volume
Al Al2O3 cBN AlBO
45 500 0 21.27 48.80 29.93 0.00
45 600 0 19.87 50.20 29.93 0.00
45 700 0 19.00 51.09 29.91 0.00
45 800 0 6.77 63.36 29.87 0.00
45 900 0 5.52 65.62 28.86 0.00
45 1000 0 3.20 69.26 27.55 0.00
45 1000 60 1.43 66.08 26.55 5.94
45 1000 180 1.34 59.56 26.24 12.85
45 1000 300 1.10 60.67 25.70 12.53
180 500 0 22.50 47.55 29.95 0.00
180 600 0 21.29 48.80 29.90 0.00
180 700 0 20.20 49.59 30.21 0.00
180 800 0 5.80 64.33 29.87 0.00
180 900 0 4.20 66.94 28.86 0.00
180 1000 0 3.16 67.39 29.44 0.00
180 1000 60 3.07 61.70 29.70 5.52
180 1000 180 2.76 58.57 28.55 10.12
180 1000 300 2.61 61.71 25.07 10.61
A.T.Shonhiwa 117
Appendix
Table A.2 Composition by volume of sample 504010 as a function of
temperature and pressure.
PRESSURE MPa
TEMP oC
TIME MINS
% Composition Volume
Al Al2O3 cBN AlBO
45 500 0 17.08 72.99 9.93 0.00
45 600 0 15.97 74.35 9.68 0.00
45 700 0 15.42 74.61 9.98 0.00
45 800 0 5.19 84.90 9.90 0.00
45 900 0 4.90 85.22 9.88 0.00
45 1000 0 3.75 86.37 9.87 0.00
45 1000 60 4.32 85.92 9.76 0.00
45 1000 180 1.76 82.66 9.69 5.89
45 1000 300 1.63 82.46 9.39 6.51
180 500 0 21.57 68.43 10.00 0.00
180 600 0 19.30 70.70 10.00 0.00
180 700 0 16.42 73.60 9.98 0.00
180 800 0 5.32 84.78 9.90 0.00
180 900 0 4.21 85.91 9.88 0.00
180 1000 0 3.60 86.54 9.86 0.00
180 1000 60 2.03 88.32 9.65 0.00
180 1000 180 1.31 83.22 9.32 6.16
180 1000 300 1.23 82.26 9.04 7.47
A.T.Shonhiwa 118
Appendix
Table A.3 Composition by volume of sample 454510 as a function of
temperature and pressure.
PRESSURE MPa
TEMP oC
TIME MINS
% Composition Volume
Al Al2O3 cBN AlBO
45 500 0 13.21 76.80 10.00 0.00
45 600 0 12.50 77.50 10.00 0.00
45 700 0 12.01 78.01 9.97 0.00
45 800 0 5.15 85.02 9.83 0.00
45 900 0 4.15 86.04 9.81 0.00
45 1000 0 3.50 86.72 9.78 0.00
45 1000 60 2.70 87.67 9.63 0.00
45 1000 180 1.14 85.22 9.49 4.14
45 1000 300 1.12 82.94 9.37 6.57
180 500 0 18.50 71.51 10.00 0.00
180 600 0 15.30 74.70 10.00 0.00
180 700 0 13.02 77.01 9.97 0.00
180 800 0 5.51 84.67 9.83 0.00
180 900 0 5.20 85.00 9.81 0.00
180 1000 0 4.05 86.60 9.35 0.00
180 1000 60 2.18 88.51 9.30 0.00
180 1000 180 1.38 82.97 9.11 6.55
180 1000 300 1.36 82.46 8.91 7.28
A.T.Shonhiwa 119
Appendix
Table A.4 Composition by volume of sample 505000 as a function of
temperature and pressure.
PRESSURE MPa
TEMP oC
TIME MINS
% Composition Volume
Al Al2O3 cBN AlBO
45 500 0 17.20 82.80 0.00 0.00
45 600 0 13.31 86.69 0.00 0.00
45 700 0 10.11 89.89 0.00 0.00
45 800 0 4.50 95.50 0.00 0.00
45 900 0 4.03 95.97 0.00 0.00
45 1000 0 2.92 97.08 0.00 0.00
45 1000 60 0.00 100.00 0.00 0.00
45 1000 180 0.00 100.00 0.00 0.00
45 1000 300 0.00 100.00 0.00 0.00
180 500 0 18.25 81.75 0.00 0.00
180 600 0 16.30 83.70 0.00 0.00
180 700 0 11.12 88.88 0.00 0.00
180 800 0 5.05 94.95 0.00 0.00
180 900 0 4.50 95.50 0.00 0.00
180 1000 0 3.48 96.52 0.00 0.00
180 1000 60 1.74 98.26 0.00 0.00
180 1000 180 1.26 98.74 0.00 0.00
180 1000 300 1.19 98.81 0.00 0.00
A.T.Shonhiwa 120
Appendix
Table A.5 Composition by mass of sample 502030 as a function of temperature and pressure.
PRESSURE
MPa TEMP
oC
TIME MINS
% Composition Mass
Al Al2O3 cBN AlBO
45 500 0 16.14 54.59 29.27 0.00
45 600 0 15.00 55.87 29.13 0.00
45 700 0 14.30 56.68 29.02 0.00
45 800 0 4.88 67.35 27.76 0.00
45 900 0 3.96 69.37 26.68 0.00
45 1000 0 2.27 72.51 25.22 0.00
45 1000 60 1.03 70.10 24.63 4.24
45 1000 180 0.99 64.68 24.92 9.40
45 1000 300 0.81 65.71 24.34 9.14
180 500 0 17.15 53.43 29.42 0.00
180 600 0 16.16 54.59 29.25 0.00
180 700 0 15.28 55.28 29.45 0.00
180 800 0 4.17 68.16 27.67 0.00
180 900 0 3.00 70.45 26.56 0.00
180 1000 0 2.25 70.73 27.02 0.00
180 1000 60 2.23 66.01 27.78 3.98
180 1000 180 2.03 63.51 27.07 7.39
180 1000 300 1.91 66.69 23.69 7.72
A.T.Shonhiwa 121
Appendix
Table A.6 Composition by mass of sample 504010 as a function of temperature and pressure.
PRESSURE
MPa TEMP
oC
TIME MINS
% Composition Mass
Al Al2O3 cBN AlBO
45 500 0 12.42 78.26 9.31 0.00
45 600 0 11.57 79.39 9.04 0.00
45 700 0 11.15 79.54 9.30 0.00
45 800 0 3.63 87.45 8.92 0.00
45 900 0 3.42 87.69 8.89 0.00
45 1000 0 2.61 88.55 8.85 0.00
45 1000 60 3.01 88.23 8.76 0.00
45 1000 180 1.24 85.84 8.80 4.12
45 1000 300 1.15 85.74 8.54 4.56
180 500 0 15.94 74.54 9.52 0.00
180 600 0 14.15 76.40 9.45 0.00
180 700 0 11.92 78.75 9.34 0.00
180 800 0 3.72 87.36 8.92 0.00
180 900 0 2.93 88.20 8.87 0.00
180 1000 0 2.50 88.66 8.83 0.00
180 1000 60 1.40 90.00 8.60 0.00
180 1000 180 0.92 86.32 8.45 4.30
180 1000 300 0.87 85.66 8.23 5.24
A.T.Shonhiwa 122
Appendix
Table A.7 Composition by mass of sample 454510 as a function of
temperature and pressure.
PRESSURE MPa
TEMP oC
TIME MINS
% Composition Mass
Al Al2O3 cBN AlBO
45 500 0 9.48 81.27 9.25 0.00
45 600 0 8.95 81.82 9.23 0.00
45 700 0 8.59 82.22 9.19 0.00
45 800 0 3.60 87.54 8.85 0.00
45 900 0 2.89 88.31 8.80 0.00
45 1000 0 2.43 88.81 8.76 0.00
45 1000 60 1.87 89.53 8.60 0.00
45 1000 180 0.80 87.78 8.55 2.87
45 1000 300 0.79 86.11 8.51 4.59
180 500 0 13.52 77.05 9.42 0.00
180 600 0 11.06 79.62 9.32 0.00
180 700 0 9.34 81.44 9.22 0.00
180 800 0 3.85 87.28 8.86 0.00
180 900 0 3.63 87.53 8.83 0.00
180 1000 0 2.82 88.79 8.38 0.00
180 1000 60 1.51 90.20 8.29 0.00
180 1000 180 0.97 86.17 8.27 4.58
180 1000 300 0.96 85.84 8.11 5.10
A.T.Shonhiwa 123
Appendix
Table A.8 Composition by mass of sample 505000 as a function of temperature and pressure.
PRESSURE
MPa TEMP
oC
TIME MINS
% Composition Mass
Al Al2O3 cBN AlBO
45 500 0 12.35 87.65 0.00 0.00
45 600 0 9.43 90.57 0.00 0.00
45 700 0 7.09 92.91 0.00 0.00
45 800 0 3.10 96.90 0.00 0.00
45 900 0 2.77 97.23 0.00 0.00
45 1000 0 2.00 98.00 0.00 0.00
45 1000 60 0.00 100.00 0.00 0.00
45 1000 180 0.00 100.00 0.00 0.00
45 1000 300 0.00 100.00 0.00 0.00
180 500 0 13.15 86.85 0.00 0.00
180 600 0 11.67 88.33 0.00 0.00
180 700 0 7.82 92.18 0.00 0.00
180 800 0 3.48 96.52 0.00 0.00
180 900 0 3.10 96.90 0.00 0.00
180 1000 0 2.39 97.61 0.00 0.00
180 1000 60 1.19 98.81 0.00 0.00
180 1000 180 0.86 99.14 0.00 0.00
180 1000 300 0.81 99.19 0.00 0.00
A.T.Shonhiwa 124
Appendix
Table A.9 Mass change as a function of temperature and pressure for
sample 502030. PRESSURE
MPa TEMP
oC
TIME MINS
W1
Grams W2
Grams
dWTh 100*12
ThdW
WW
45 500 0 0.9741 1.0706 37.24 26.60
45 600 0 0.9744 1.1001 37.24 34.64
45 700 0 0.9825
1.1488
37.24 45.40
45 800 0 0.9645 1.1366 37.24 47.91
45 900 0 1.031 1.2495 37.24 56.91
45 1000 0 0.9918 1.2299 37.24 64.47
45 1000 60 0.8871 1.1906 37.24 91.87
45 1000 180 0.7344 0.9870 37.24 92.36
45 1000 300 0.5875 0.7899 37.24 92.51
180 500 0 1.0113 1.1388 37.24 33.85
180 600 0 0.9711 1.1000 37.24 35.64
180 700 0 0.9925 1.1503 37.24 42.69
180 800 0 0.9745
1.1402
37.24 45.69
180 900 0 1.061 1.2795 37.24 55.30
180 1000 0 1.0099 1.2390 37.24 60.92
180 1000 60 0.8933 1.1900 37.24 89.19
180 1000 180 0.8026 1.0750 37.24 91.14
180 1000 300 0.8664 1.1600 37.24 91.00
A.T.Shonhiwa 125
Appendix
Table A.10 Mass change as a function of temperature and pressure for sample 504010. PRESSURE
MPa TEMP
oC
TIME MINS
W1
Grams
W2
Grams
dWTh 100*12
ThdW
WW
45 500 0 0.8512 0.9326 36.11 26.48
45 600 0 0.8535 0.9433 36.11 29.14
45 700 0 0.8826 1.0050 36.11 38.41
45 800 0 0.8454 0.9908 36.11 47.63
45 900 0 0.8705 1.0207 36.11 47.78
45 1000 0 0.8722 1.0230 36.11 47.88
45 1000 60 0.9002 1.1710 36.11 83.31
45 1000 180 0.768 1.0010 36.11 84.02
45 1000 300 0.776 1.0290 36.11 90.29
180 500 0 0.8835
0.9685
36.11 29.68
180 600 0 0.873 0.9579 36.11 26.97
180 700 0 0.8826 0.9953 36.11 35.50
180 800 0 0.8454
0.9803
36.11 44.37
180 900 0 0.8705 1.0205 36.11 47.72
180 1000 0 0.8722 1.0220 36.11 47.56
180 1000 60 1.008 1.2900 36.11 77.47
180 1000 180 0.6367 0.8253 36.11 82.03
180 1000 300 0.7618 0.9875 36.11 82.05
A.T.Shonhiwa 126
Appendix
Table A.11 Mass change as a function of temperature and pressure for
sample 454510. PRESSURE
MPa TEMP
oC
TIME MINS
W1
Grams W2
Grams
dWTh 100*12
ThdW
WW
45 500 0 0.9337 1.0258 31.88 30.94
45 600 0 0.9253 1.0308 31.88 35.76
45 700 0 0.9248 1.0620 31.88 46.54
45 800 0 0.9256 1.0762 31.88 51.04
45 900 0 0.9363 1.1033 31.88 55.95
45 1000 0 0.9719 1.1463 31.88 56.29
45 1000 60 1.0082 1.2606 31.88 78.53
45 1000 180 0.926 1.1625 31.88 80.11
45 1000 300 0.8994 1.1340 31.88 81.82
180 500 0 0.95 1.049 31.88 32.66
180 600 0 0.9326 1.0439 31.88 37.53
180 700 0 0.9248
1.05
31.88 42.64
180 800 0 0.9256 1.0850 31.88 54.02
180 900 0 0.9363 1.1030 31.88 55.85
180 1000 0 0.9719 1.1463 31.88 56.29
180 1000 60 0.3599 0.4475 31.88 76.35
180 1000 180 0.8532 1.0672 31.88 78.68
180 1000 300 1.0116 1.2655 31.88 78.73
A.T.Shonhiwa 127
Appendix
Table A.12 Mass change as a function of temperature and pressure for
sample 505000. PRESSURE
MPa TEMP
oC
TIME MINS
W1
Grams W2
Grams
dWTh 100*12
ThdW
WW
45 500 0 0.6901 0.7695 35.57 32.35
45 600 0 0.6909 0.8031 35.57 45.66
45 700 0 0.6639 0.7831 35.57 50.48
45 800 0 0.6479 0.7708 35.57 53.33
45 900 0 0.6904 0.8264 35.57 55.38
45 1000 0 0.688 0.8355 35.57 60.29
45 1000 60 0.6935 0.9136 35.57 89.23
45 1000 180 0.7348 0.9746 35.57 91.75
45 1000 300 0.7283 0.9653 35.57 91.49
180 500 0 0.6793 0.7648 35.57 35.38
180 600 0 0.6899 0.8000 35.57 44.87
180 700 0 0.6639
0.7781
35.57 48.35
180 800 0 0.6479
0.7636
35.57 50.20
180 900 0 0.6904 0.8260 35.57 55.22
180 1000 0 0.5323 0.6933 35.57 85.03
180 1000 60 0.5327 0.6934 35.57 84.81
180 1000 180 0.7113 0.9365 35.57 89.01
180 1000 300 0.7432 0.9845 35.57 91.28
A.T.Shonhiwa 128
Appendix
Table A.13 Degree of reaction as a function of temperature and
pressure for sample 502030. PRESSURE
MPa TEMP
oC
TIME MINS
Al Initial
(g)
Al final (g) 100*1
initial
final
Al
Al
45 500 0 0.412 0.17 58.08
45 600 0 0.412 0.17 59.98
45 700 0 0.416 0.16 61.12
45 800 0 0.408 0.06 86.41
45 900 0 0.436 0.05 88.66
45 1000 0 0.420 0.03 93.34
45 1000 60 0.375 0.01 96.72
45 1000 180 0.311 0.01 96.85
45 1000 300 0.249 0.01 97.42
180 500 0 0.428 0.20 54.37
180 600 0 0.411 0.18 56.73
180 700 0 0.420 0.18 58.16
180 800 0 0.412 0.05 88.11
180 900 0 0.449 0.04 91.46
180 1000 0 0.427 0.03 93.48
180 1000 60 0.378 0.03 92.98
180 1000 180 0.340 0.02 93.57
180 1000 300 0.367 0.02 93.97
A.T.Shonhiwa 129
Appendix
Table A.14 Degree of reaction as a function of temperature and pressure for sample 504010.
PRESSURE
MPa TEMP
oC
TIME MINS
Al Initial
(g)
Al final (g) 100*1
initial
final
Al
Al
45 500 0 0.349 0.12 66.82
45 600 0 0.350 0.11 68.83
45 700 0 0.362 0.11 69.05
45 800 0 0.347 0.04 89.62
45 900 0 0.357 0.03 90.22
45 1000 0 0.358 0.03 92.55
45 1000 60 0.369 0.04 93.46
45 1000 180 0.315 0.01 96.07
45 1000 300 0.318 0.01 96.27
180 500 0 0.363 0.16 55.59
180 600 0 0.358 0.14 60.54
180 700 0 0.362 0.12 65.97
180 800 0 0.347 0.04 89.38
180 900 0 0.357 0.03 91.62
180 1000 0 0.358 0.03 92.85
180 1000 60 0.414 0.02 95.62
180 1000 180 0.261 0.01 97.09
180 1000 300 0.313 0.01 97.24
A.T.Shonhiwa 130
Appendix
Table A.15 Degree of reaction as a function of temperature and pressure for sample 454510.
PRESSURE
MPa TEMP
oC
TIME MINS
Al Initial
(g)
Al final (g) 100*1
initial
final
Al
Al
45 500 0 0.383 0.10 74.63
45 600 0 0.380 0.09 75.69
45 700 0 0.335 0.09 72.76
45 800 0 0.335 0.04 88.43
45 900 0 0.339 0.03 90.60
45 1000 0 0.352 0.03 92.08
45 1000 60 0.365 0.02 93.54
45 1000 180 0.335 0.01 97.24
45 1000 300 0.326 0.01 97.24
180 500 0 0.344 0.15 57.59
180 600 0 0.338 0.12 64.84
180 700 0 0.335 0.10 69.89
180 800 0 0.335 0.04 87.53
180 900 0 0.339 0.04 88.19
180 1000 0 0.352 0.03 90.81
180 1000 60 0.130 0.01 94.80
180 1000 180 0.309 0.01 96.65
180 1000 300 0.367 0.01 96.69
A.T.Shonhiwa 131
Appendix
Table A.16 Degree of reaction as a function of temperature and pressure for sample 505000.
PRESSURE
MPa TEMP
oC
TIME MINS
Al Initial
(g)
Al final (g) 100*1
initial
final
Al
Al
45 500 0 0.250 0.10 74.90
45 600 0 0.250 0.08 69.76
45 700 0 0.268 0.06 79.31
45 800 0 0.262 0.02 90.88
45 900 0 0.279 0.02 91.80
45 1000 0 0.278 0.02 94.08
45 1000 60 0.280 0.00 100.00
45 1000 180 0.297 0.00 100.00
45 1000 300 0.294 0.00 100.00
180 500 0 0.275 0.10 62.55
180 600 0 0.279 0.09 66.52
180 700 0 0.268 0.06 77.20
180 800 0 0.262 0.03 89.77
180 900 0 0.279 0.03 90.83
180 1000 0 0.215 0.02 92.31
180 1000 60 0.215 0.01 96.16
180 1000 180 0.288 0.01 97.20
180 1000 300 0.300 0.01 97.34
A.T.Shonhiwa 132
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